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Title:
AQUEOUS ELETRODEPOSITION OF MAGNETIC COBALT-SAMARIUM ALLOYS
Kind Code:
A1
Abstract:
Disclosed are methods and compositions for aqueous electrodeposition of rare earth-transitiona metal alloys (e.g., samarium-cobalt alloys). Also disclosed are nanostructured magnetic coatings comprising a magnetic alloy of a rare earth metal (e.g., samarium) and a transition metal (e.g., cobalt). This abstract is intended as a scanning tool for purposes of searching in the particular art and is not intended to be limiting of the present invention.


Inventors:
Nobe, Ken (US)
Application Number:
11/872699
Publication Date:
10/02/2008
Filing Date:
10/15/2007
Primary Class:
Other Classes:
205/238
International Classes:
C25D3/56
View Patent Images:
Attorney, Agent or Firm:
Ballard Spahr Andrews & Ingersoll, LLP (SUITE 1000, 999 PEACHTREE STREET, ATLANTA, GA, 30309-3915, US)
Claims:
What is claimed is:

1. A composition for enhancing the aqueous electrodeposition of rare earth-transition metal alloys comprising: a water soluble salt of samarium, a water soluble salt of cobalt, and a complexant.

2. The composition of claim 1, wherein the water soluble salt of samarium is samarium sulfamate.

3. The composition of claim 1, wherein the water soluble salt of cobalt is cobalt sulfate or cobalt sulfamate.

4. The composition of claim 1, wherein the complexant is selected from one or more amine carboxylates, one or more hydroxycarboxylic acids, and combinations thereof.

5. The composition of claim 1, further comprising one or more supporting electrolytes.

6. The composition of claim 5, wherein the one or more supporting electrolytes is selected from ammonium sulfamate, ammounium sulfate, ammonium chloride, and mixtures thereof.

7. The composition of claim 1, comprising from about 0.25M to about 2.0M of the water soluble salt of samarium, from about 0.01M to about 0.5M of the water soluble salt of cobalt, from about 0.05M to about 0.5M of the complexant, and from about 0M to about 3M of one or more supporting electrolytes.

8. The composition of claim 7, comprising about 1M of the water soluble salt of samarium, about 0.05M of the water soluble salt of cobalt, about 0.15M of the complexant, and about 1M of the one or more supporting electrolytes.

9. A method for electrodepositing a samarium-cobalt coating onto a conducting substrate, comprising: a. placing an aqueous solution containing a water soluble salt of samarium, a water soluble salt of cobalt, one or more supporting electrolytes, and a complexant into a plating bath, b. placing an anode and the substrate to be coated into the bath and connecting the anode and the substrate to a power supply, with the substrate acting as a cathode, c. adjusting the pH of the bath to a suitable operating level, and d. applying a current through the anode and substrate causing the samarium and the cobalt to migrate to, and adhere to, the substrate.

10. The method of claim 9, wherein the water soluble salt of samarium is samarium sulfamate.

11. The method of claim 9, wherein the water soluble salt of cobalt is cobalt sulfate or cobalt sulfamate.

12. The method of claim 9, wherein the complexant is selected from one or more amine carboxylates, one or more hydroxycarboxylic acids, and combinations thereof.

13. The method of claim 9, wherein the one or more supporting electrolytes is selected from ammonium sulfamate, ammounium sulfate, ammonium chloride, and mixtures thereof.

14. The method of claim 9, wherein the aqueous solution further comprises boric acid.

15. The method of claim 9, wherein the aqueous solution comprises from about 0.25M to about 2.0M of the water soluble salt of samarium, from about 0.01M to about 0.5M of the water soluble salt of cobalt, from about 0.05M to about 0.5M of the complexant, and from about 0.0001M to about 3M of the supporting electrolytes.

16. The method of claim 15, wherein the aqueous solution comprises about 1M of the water soluble salt of samarium, about 0.05M of the water soluble salt of cobalt, about 0.15M of the complexant, and about 1M of the supporting electrolytes.

17. The method of claim 9, wherein a current density of from about 5 mA/cm2 to about 600 mA/cm2 is applied across the anode and cathode.

18. The method of claim 9, wherein the current is applied with pulse current modifications varying with duty cycle and frequency.

19. The method of claim 9, wherein the pH of the solution is from about 4 to about 6.5.

20. The method of claim 9, wherein the solution temperature is adjusted to from about 25° C. to about 60° C.

21. A samarium-cobalt coating produced by the method of claim 9.

22. A nanostructured magnetic coating comprising a magnetic alloy of a rare earth metal and a transition metal.

23. The nanostructured magnetic coating of claim 22, wherein the rare earth metal is samarium and wherein the transition metal is cobalt.

24. The nanostructured magnetic coating of claim 22, wherein the coating is provided by electrodeposition from an aqueous solution.

25. The nanostructured magnetic coating of claim 22, wherein the alloy comprises SmCo5 or Sm2Co17.

Description:

CROSS-REFERENCE TO RELATED APPLICATIONS

This application claims the benefit of U.S. Application No. 60/851,389, filed Oct. 13, 2006, and U.S. Application No. 60/852,286, filed Oct. 17, 2006, which are hereby incorporated herein by reference in their entireties.

ACKNOWLEDGEMENT

This invention was made with government support under Grant No. DMI-0089095 awarded by the National Science Foundation. The United States government has certain rights in the invention.

BACKGROUND

Bulk alloys of transition metals-rare earths are important permanent magnet materials. When using conventional techniques, however, the current high materials and processing costs of Co—Sm permanent magnets have limited their application to high temperature and corrosive environments where costs are of secondary importance. More specifically, only high cost metallurgical and physical deposition methods are currently in use to fabricate Co—Sm permanent magnets consisting of the intermetallics SmCo5 and Sm2Co17.

In contrast, compositions and methods disclosed in U.S. Pat. No. 6,306,276 established the basis for the successful electrodeposition of rare earth-transition metal alloys from aqueous media. Suitable operating and plating bath conditions to obtain magnetic cobalt-samarium (Co-Sm) alloys from aqueous media for high performance nanostructured permanent magnets, however, have remained unknown in the art.

As current estimates of the global market for permanent magnets exceed $5 billion, such suitable operating and plating bath conditions to obtain magnetic cobalt-samarium (Co—Sm) alloys can provide substantial savings in manufacturing costs and considerable lower materials costs for nanotechnology applications, thereby greatly expanding the global market share of high performance Co—Sm permanent magnets fabricated by electrodeposition from aqueous media.

SUMMARY

As embodied and broadly described herein, the invention, in one aspect, relates to

Disclosed are compositions for enhancing the aqueous electrodeposition of rare earth-transition metal alloys comprising a water soluble salt of samarium, a water soluble salt of cobalt, and a complexant.

Also disclosed are methods for electrodepositing a samarium-cobalt coating onto a conducting (e.g., metal) substrate, comprising placing an aqueous solution containing a water soluble salt of samarium, a water soluble salt of cobalt, one or more supporting electrolytes, and a complexant into a plating bath, placing an anode and the substrate to be coated into the bath and connecting the anode and the substrate to a power supply, with the substrate acting as a cathode, adjusting the pH of the bath to a suitable operating level, and applying a current through the anode and substrate causing the samarium and the cobalt to migrate to, and adhere to, the substrate.

Also disclosed are samarium-cobalt coatings produced by the disclosed methods.

Also disclosed are nanostructured magnetic coatings comprising a magnetic alloy of a rare earth metal and a transition metal.

Unless otherwise expressly stated, it is in no way intended that any method or aspect set forth herein be construed as requiring that its steps be performed in a specific order. Accordingly, where a disclosed method or system does not specifically state that the steps are to be limited to a specific order, it is no way intended that an order be inferred, in any respect. This holds for any possible non-express basis for interpretation, including matters of logic with respect to arrangement of steps or operational flow, plain meaning derived from grammatical organization or punctuation, or the number or type of aspects described in the specification.

BRIEF DESCRIPTION OF THE FIGURES

The accompanying figures, which are incorporated in and constitute a part of this specification, illustrate several aspects and together with the description serve to explain the principles of the invention.

FIG. 1 shows a graph illustrating maximum energy product for permanent magnetic materials as a function of time.

FIG. 2 shows a phase diagram of the Sm—Co system.

FIG. 3 shows the Co-rich region of the Sm—Co phase diagram.

FIG. 4 shows the hexagonal CaCu5 structure of SmCo5 [C. Barrett and T. B. Massalski, Structure of metals, Pergamon Press, Oxford, New York, (1980), p266].

FIG. 5 shows the rhombohedral Th2Ni17-type structure of Sm2Co17 [R. C. O'Handley, Modern magnetic materials, John Wiley & Son, Inc., New York, (2000), pp. 496-502].

FIG. 6 shows the pinning mechanism of Sm2Co17 [P. Campbell, Permanent magnet materials and their application, Cambridge University Press, New York, (1994), pp. 42-43].

FIG. 7 shows saturation magnetization, coercivity and squareness of Co—Sm films versus atomic percent of Sm [S. A. Bendson and J. H. Judy, IEEE Trans. Magnetics, 9, 627, (1973)].

FIG. 8 shows variation of resistivity, coercivity, and saturation magnetization with film (33 nm) composition (left) and SUBSTRATE temperature (right) under a background pressure of 2×10−7 Torres.

FIG. 10 shows coercivity (left) and remanence ratio (right) as a function of the Sm content for samples processed at Ar gas pressures ▪6×10−2, 8×10−2, and ♦1×10−1 Torr and a constant substrate temperature 460° C. [V. Neua and S. A. Shaheen, J. Appl. Phys., 86, 7006, (1999)].

FIG. 11 is of graphs showing the effect of current density, with oscillatory stirring, on the co-deposition of rare earth TM alloyed with nickel, iron and cobalt respectively.

FIG. 12 is a graph showing the effect, with stirring, of glycine/cobalt ratio on the deposition of the rare earth cobalt mixture.

FIG. 13 is a graph showing the effect, with stirring, of glycine and cobalt concentration on rare earth cobalt mixture deposition.

FIG. 14 is a graph showing the effect, with stirring, of pulse current duty cycle on rare earth cobalt mixture deposition.

FIG. 15 is of graphs showing the effect of solution pH and current density on the deposition of Nd—Ni, Nd—FE and Nd—Co, respectively.

FIG. 16 is of graph showing the effect of solution stirring on Ce—Ni deposits.

FIG. 17 shows an experimental flowchart of a Hull cell study of DC and PC electrodeposition of Co—Sm alloys.

FIG. 18 shows a design of Hull cell and Hull cell panel.

FIG. 19 shows a schematic of pulse current electrodeposition.

FIG. 20 shows a schematic of the setup for the Hull cell electrodeposition system.

FIG. 21 shows an energy dispersive spectrum of an electrodeposited Co—Sm alloy.

FIG. 22 shows Hull cell patterns of PC electrodeposition from bath 1 at 60° C. for different applied current of (a) 4.5 A and (b) 7 A. (Ton=10 ms, duty cycle γ=0.1 and applied charge=50 C).

FIG. 23 shows Hull cell patterns of PC electrodeposition from bath I at 25° C. for applied charge/deposit area of (a) 100 C/15 cm2 and (b) 50 C/7.5 cm2. (Ton=10 ms, duty cycle γ=0.1 and an applied current of 4.5 A).

FIG. 24 shows Hull cell patterns obtained from bath 1 at (a) 25° C., (b) 60° C. and (c) 80° C.

FIG. 25 shows XRD patterns of deposits #3 obtained from bath 1 at 25° C. and (a) 50 mA/cm2, metallic region; (b) 100 mA/cm2, burnt region; and (c) 500 mA/cm2, oxide/hydroxide region.

FIG. 26 shows Sm deposit content obtained from bath 1 at 25, 60 and 80° C. and various CDs.

FIG. 27 shows a setup for solution agitation in the Hull cell.

FIG. 28 shows Hull cell patterns obtained from Bath 1 ar 25° C. (a) without and (b) with agitation; at 60° C. (c) without and (d) with agitation.

FIG. 29 shows Sm deposit content obtained from bath 1 with/out agitation at 25 and 60° C. and various CD.

FIG. 30 shows Hull cell patterns obtained from Bath 2 (1M Sm sulfamate without glycine) at (a) 25° C. and (c) 60° C.; from bath 3 (1M Sm sulfamate with 0.15M glycine) at (b) 25° C. and (d) 60° C.

FIG. 31 shows Hull cell patterns obtained from Bath 4 (0.05M Co sulfate) at (a) 25° C. and (b) 60° C.; from bath 5 (0.05M Co sulfate, 0.15M glycine) at (c) 25° C. and (d) 60° C.

FIG. 32 shows Hull cell patterns obtained from bath 6 (1M Sm sulfamate, 0.05M Co sulfate), bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine) and bath 7 (1M Sm sulfamate, 0.05M Co sulfate, 3M glycine) at (a)-(c) 25° C. and (d)-(t) 60° C.

FIG. 33 shows Sm deposit content obtained from plating baths containing no, 0.15M and 3M glycine at various CD.

FIG. 34 shows XRD patterns of metallic deposits obtained from bath 1 (with 0.15M glycine) at (a) 60° C., 650 mA/cm2and (b) 25° C., 50 mA/cm2 and from bath 6 (without glycine) at (c) 60° C., 50 mA/cm2 and (d) 25° C., 10 mA/cm2.

FIG. 35 shows Hull cell patterns obtained from Bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine) at (a) 25° C. and (b) 60° C.; from Bath 8 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, 1M NH4 sulfamate) at (c) 25° C. and (d) 60° C.

FIG. 36 shows Sm deposit content obtained at 25 and 60° C. from Bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine) and Bath 8 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, 1M sulfamate) at various CD.

FIG. 37 shows Hull cell patterns obtained from Bath 1 at 25° C. for (a) DC, (b) γ=0.1, T.=10 ms and (c) γ=0.1, Ton=0.05 ms; at 60° C. for (d) DC, (e) γ=0.1, Ton=10 ms and (f) γ=0.1, T.=0.05 ms.

FIG. 38 shows Sm deposit content of deposits obtained from Bath 1 at 25 and 60° C. for PC electrodeposition of LH=0.05 and 10 ms (γ=0.1) and for DC electrodeposition at various PCD.

FIG. 39 shows Hull cell patterns obtained from Bath 1 at 25° C. for γ=0.1 (a) Ton=0.05 ms, (b) Ton=0.1 ms, (c) Ton=1 ms and (d) Ton=10 ms.

FIG. 40 shows Sm deposit content of deposits obtained from Bath 1 at 25° C. and γ=0.1 for Ton=0.05, 0.1, 1 and 10 ms at various PCD.

FIG. 41 shows Hull cell patterns obtained from Bath 1 at 25° C. for Ton=0.1 ms at (a) γ=0.05, (b) γ=0.075, (c) γ=0.1, (d) γ=0.2, (e) γ=0.3 and (f) γ=0.3 (DC).

FIG. 42 shows Sm deposit content obtained from Bath 1 at 25° C. and Ton=0.1 ms for γ=0.05, 0.075, 0.1, 0.2, 0.3 and 1 (DC) at various PCD.

FIG. 43 shows dependence of deposit Sm content on current density.

FIG. 44 shows dependence of current efficiency on current density.

FIG. 45 shows dependence of deposit Sm content on temperature.

FIG. 46 shows dependence of current efficiency on temperature.

FIG. 47 shows dependence of magnetic saturation on deposit Sm content (no NH4 sulfamate).

FIG. 48 shows dependence of coercivity on current density.

FIG. 49 shows topography and microstructure of electrodeposited Co—Sm alloys at current densities 100 mA/cm2 for plating baths of 60° C. (a) with 1M NH4 sulfamate, and (b) without NH4 sulfamate.

FIG. 50 shows XRD of electrodeposited Co—Sm alloys at 100 and 500 mA/cm2. S=brass substrate.

FIG. 51 shows the effect of CD on Sm content and CE of Co—Sm alloys. 1 M Sm sulfamate, 0.05 M Co sulfate, 0.15 M.

FIG. 52 shows the effect of CD on deposit composition; pH 6, 1.2 p.m film thickness.

FIG. 53 shows structures of IG-RE-Glycine complexes (17): (a) equilibrium of anionic, zwitterionic and cationic species; (b) hetero-dinuclear trisglycine complex; (c) quasi-diglycine complex; (d) quasi-triglycine complex.

FIG. 54 shows structures of Co—V (a) and Co—Fe—V (b) biscitrate complexes.

FIG. 55 shows proposed mechanism of electrodeposition of binary and ternary IG-V alloys.

FIG. 56 shows structure of Co—Mo (W) biscitrate complexes.

FIG. 57 shows a proposed mechanism of electrodeposition of IG-Mo (W) alloys.

FIG. 58 shows composite Co—W/Cr/Co—W/Cr deposit (12, 24): (a) X500, not heat treated (HT); (b) H. T. in air, 916° C., 10 hrs (unetched); (c) H. T. in carburizing atmosphere, 916° C., 10 hrs (unetched); (d) H. T. in carburizing atmosphere, 916° C., 10 hrs (etchant, hot Murakami). S=cobalt strike.

FIG. 59 shows an experimental flowchart of parametric studies of DC electrodeposition of Co—Sm alloys.

FIG. 60 shows a set up of DC electrodeposition.

FIG. 61 shows a setup of a RDE system.

FIG. 62 shows a schematic diagram of a VSM.

FIG. 63 shows a hysteresis loop of ED Co—Sm alloys.

FIG. 64 shows the effect of current density and solution temperature on (a) samarium deposit content and (b) current efficiency.

FIG. 65 shows the effect of current density and solution temperature on normalized charges of Sm, Co and H2 in Co—Sm alloys electrodeposition.

FIG. 66 shows (a) polarization curves at various current densities and solution temperatures and (b) dependence of Sm content on cathodic potential in Co—Sm electrodeposition.

FIG. 67 shows XRD patterns of deposits obtained from bath 1 (1M Sm sulfamate, 0.051M Co sulfate, 0.15M glycine, pH 6) at 25° C. and various CDs).

FIG. 68 shows XRD patterns of deposits obtained from bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, pH 6) at 60° C. and various CDs.

FIG. 69 shows XRD patterns of deposits obtained from bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, pH 6) at (a) 2 mA/cm2, (b) 25 mA/cm2 and (c) 50 mA/cm2 at various solution temperatures.

FIG. 70 shows SEM of Co—Sm thin films obtained from bath 1 at 25° C. and at (a)-(c) 2 mA/cm2, (d)-(f) 25 mA/cm2 and (g)-(i) 50 mA/cm2.

FIG. 71 shows SEM of Co—Sm thin films obtained from bath 1 at 60° C. and at (a) (c) 25 mA/cm2, (d)-(f) 50 mA/cm2, (g)-(i) 100 mA/cm2, (j)-(l) 300 mA/cm2 and (m)-(o) 500 mA/cm2.

FIG. 72 shows SEM of Co—Sm thin films obtained from bath 1 at 50 mA/cm2 and at (a)-(b): 25° C., (c)-(d) 40° C. and (e)-(f) 60° C.

FIG. 73 shows SEM of Co—Sm thin films obtained from bath 1 at 50 mA/cm2 and at (a)-(c) 25° C., (d)-(f) 40° C. and (g)-(i) 60° C.

FIG. 74 shows dependence of particle size on Sm deposit content at various temperatures and CDs.

FIG. 75 shows magnetic hysteresis loops obtained at (a)-(c) 25° C. and (d)-(i) 60° C. and at various CDs from bath 1.

FIG. 76 shows effects of current density and temperature on deposit crystalline structures, particle sizes and magnetic properties.

FIG. 77 shows the effect of the particle size on coercivities of fiber-shaped microstructures in Co—Sm alloys.

FIG. 78 shows effect of solution pH on (a) Sm deposit content and (b) current efficiency at 25, 60° C. and various CDs.

FIG. 79 shows XRD patterns of deposits obtained at 10 mA/cm2, 25° C. and various solution pHs. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine).

FIG. 80 shows XRD patterns of deposits obtained at 50 mA/cm2, 25° C. and various solution pHs. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine).

FIG. 81 shows XRD patterns of deposits obtained at 10 mA/cm2, 60° C. and various solution pHs. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine).

FIG. 82 shows XRD patterns of deposits obtained at 50 mA/cm2, 60° C. and various solution plls. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine).

FIG. 83 shows XRD patterns of deposits obtained at 100 mA/cm2, 60° C. and various solution pHs. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine).

FIG. 84 shows XRD patterns of deposits obtained at 300 mA/cm2, 60° C. and various solution pHs. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine).

FIG. 85 shows low magnitude (2,000×) SEM of Co—Sm thin films obtained at (a)-(c) pH 6, (d)-(f) pH 4 and (g)-(i) pH 2 at 25° C. and various CDs. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine)

FIG. 86 shows high magnitude (50,000×) SEM of Co—Sm thin films obtained at (a)-(c) pH 6, (d)-(f) pH 4 and (g)-(i) pH 2 at 25° C. and various CDs. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine).

FIG. 87 shows low magnitude (2,000×) SEM of Co—Sm thin films at (a)-(c) pH 6, (d)-(f) pH 4 and (g)-(i) pH 2 at 60° C. and various CDs. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine).

FIG. 88 shows high magnitude (50,000×) SEM of Co—Sm thin films obtained at (a)-(c) pH 6, (d)-(f) pH 4 and (g)-(i) pH 2 at 60° C. and various CDs. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine).

FIG. 89 shows dependence of particle size on Sm deposit content at various pHs, temperatures and CDs.

FIG. 90 shows magnetic properties of deposits obtained at 25, 60° C. and various pHs. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine).

FIG. 91 shows the effect of rotating rate of RDE on (a) samarium deposit content and (b) current efficiency at 25° C. (Deposits of 0 rpm rotation rate were obtained from parallel electrode because poor deposits were obtained by RDE at 0 rpm.).

FIG. 92 shows the effect of rotating rate of RDE on normalized charges of Sm, Co and H2 in Co—Sm alloys electrodeposition. (Deposits of 0 rpm rotation rate were obtained from parallel electrode because poor deposits were obtained by RDE at 0 rpm.).

FIG. 93 shows SEM of Co—Sm thin films obtained from bath 1 at 100 mA/cm2, 25° C. and (a)-(c): no-agitation (non-metallic, obtained from parallel electrode), (d)-(f) 1000 rpm (Sm=12.5 at %), and (g)-(i) 200 rpm (Sm=7.2 at %).

FIG. 94 shows magnetic properties of deposits obtained at 25° C. and various rotating rates.

FIG. 95 shows the effect of Sm sulfamate concentration on (a) samarium deposit content and (b) current efficiency at 25 and 60° C.

FIG. 96 shows the effect of Sm sulfamate concentration on normalized charges of Sm, Co and H2 in Co—Sm alloy electrodeposition.

FIG. 97 shows XRD patterns of deposits obtained at 25° C., 25 mA/cm2 and various Sm sulfamate concentrations. (Bath: 0.05M Co sulfate, 0.15M glycine, Sm sulfamate varied from 0.25 to 1M, pH 6).

FIG. 98 shows XRD patterns of deposits obtained at 25° C., 50 mA/cm2 and various Sm sulfamate concentrations. (Bath: 0.05M Co sulfate, 0.15M glycine, Sm sulfamate varied from 0.25 to 1M, pH 6).

FIG. 99 shows XRD patterns of deposits obtained at 60° C., 50 mA/cm2 and various Sm sulfamate concentrations. (Bath: 0.05M Co sulfate, 0.15M glycine, Sm sulfamate varied from 0.25 to 1M, pH 6).

FIG. 100 shows XRD patterns of deposits obtained at 60° C., 100 mA/cm2 and various Sm sulfamate concentrations. (Bath: 0.05M Co sulfate, 0.15M glycine, 0.25 to 1M Sm sulfamate, pH 6).

FIG. 101 shows the effect of Sm sulfamate concentration on magnetic properties of deposits obtained at 25 and 60° C. (Bath: 0.05M Co sulfate, 0.15M glycine, Sm sulfamate varied from 0.25 to 1M, pH 6).

FIG. 102 shows the effect of glycine concentration on (a) samarium deposit content and (b) current efficiency at 25 and 60° C.

FIG. 103 shows XRD patterns of deposits obtained at 25° C., 50 mA/cm2 and various glycine concentrations. (Bath: 1M Sm sulfamate, 0.05M cobalt sulfate, glycine varied from 0 to 0.5M, pH 6).

FIG. 104 shows XRD patterns of deposits obtained at 60° C., 50 mA/cm2 various glycine concentrations. (Bath: 1M Sm sulfamate, 0.05M cobalt sulfate, glycine varied from 0 to 0.5M, pH 6).

FIG. 105 shows the effect of glycine concentration on magnetic properties of electrodeposited Co—Sm thin films obtained at 25 and 60° C.

FIG. 106 shows the effect of NH4 sulfamate concentration on (a) samarium deposit content and (b) currentefficiency at 25 and 60° C.

FIG. 107 shows the effect of NH4 sulfamate concentration on normalized charges of Sm, Co and H2 in Co—Sm alloy electrodeposition.

FIG. 108 shows XRD patterns of deposits obtained from (a) bath 8 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, 1M NH4 sulfamate, pH 5.2) and (b) bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, pH 5.7) at 25° C. and various CDs.

FIG. 109 shows XRD patterns of deposits obtained from (a) bath 8 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, 1M NH4 sulfamate, pH 5.2) and (b) bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, pH 5.7) at 60° C. and various CDs.

FIG. 110 shows SEM (2,000×) of Co—Sm thin films from bath (a)-(b) without (bath 1) and (c)-(d) with 1M NH4 sulfamate (bath 8).

FIG. 111 shows SEM (50,000×) of Co—Sm thin films from bath (a)-(c) without (bath 1) and (d)-(l) with 1M NH4 sulfamate (bath 8).

FIG. 112 shows the effect of NH4 sulfamate concentration on magnetic properties of electrodeposited Co—Sm thin films obtained at 25 and 60° C. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, NH4 sulfamate varied from 0 to 1M).

FIG. 113 shows the effect of supporting electrolyte on samarium deposit at (a) 25 and (b) 60° C.

FIG. 114 shows the effect of supporting electrolyte on current efficiency at (a) 25 and (b) 60° C.

FIG. 115 shows the effect of supporting electrolyte on normalized charge of Sm, Co and H2 at 25° C. and 60° C. in Co—Sm electrodeposition.

FIG. 116 shows SEM of Co—Sm thin films obtained from bath with (a)-(c) no supporting electrolyte, (d)-(f) 1M NH4 sulfamate, (g)-(i) 1M NH4Cl and (j)-(l) 1M KCl at 25° C. and 25 mA/cm2. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine with different types of supporting electrolytes) 60° C., 300 mA/cm2.

FIG. 117 shows SEM of Co—Sm thin films obtained from bath with (a)-(c) no supporting electrolyte, (d)-(0.1M NH4 sulfamate, (g)-(i) 1M NH4C1 and (j)-(l) 1M KCl at 60° C. and 300 mA/cm2. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine with different types of supporting electrolytes).

FIG. 118 shows the effect of types of supporting electrolytes on magnetic properties of electrodeposited Co—Sm thin films obtained at 25 and 60° C. (Bath: 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine with different types of supporting electrolytes).

FIG. 119 shows dependence of Bragg angle (2%) of (10.0) and (00.2) planes on Sm deposit content obtained at 60° C.

FIG. 120 shows an experimental flowchart of parametric studies of PC electrodeposition of Co—Sm alloys.

FIG. 121 shows a schematic of pulse current electrodeposition.

FIG. 122 shows a setup for PC electrodeposition.

FIG. 123 shows the effect of peak current density and solution temperature on (a) samarium deposit content and (b) current efficiency for DC and PC (Ton=0.1 ms, γ=0.1) electrodeposition.

FIG. 124 shows the effect of peak current density and solution temperature on normalized charges of Sm, Co and I-12 in Co—Sm alloys electrodeposition by DC and PC (Ton=0.1 ms, γ=0.1) electrodeposition.

FIG. 125 shows XRD patterns of deposits obtained from bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, pH 6) at 25 and 60° C. and various PCDs.

FIG. 126 shows SEM of Co—Sm thin films obtained from bath 1 at 25° C. and various PCDs.

FIG. 127 shows SEM of Co—Sm thin films obtained from bath 1 at 60° C. and various PCDs.

FIG. 128 shows magnetic hysteresis loops obtained at (a)-(c) 25° C. and (d)-(i) 60° C. and at various PCDs from bath 1.

FIG. 129 shows effects of peak current density and temperature on magnetic properties.

FIG. 130 shows the effect of duty cycle on (a) samarium deposit content and (b) current efficiency (Ton=0.1 ms).

FIG. 131 shows the effect of duty cycle on normalized charges of Sm, Co and H2, in Co—Sm alloys electrodeposition (Ton=0.1 ms, γ=0.1).

FIG. 132 shows XRD patterns of deposits obtained from bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, pH 6) at 25 and 60° C., 500 mA/cm2 and various duty cycles.

FIG. 133 shows SEM of Co—Sm thin films obtained from bath 1 at 25 and 60° C., 500 mA/cm2, and various duty cycles.

FIG. 134 shows the effects of duty cycles on magnetic properties.

FIG. 135 shows the effect of frequency on (a) samarium deposit content and (b) current efficiency (γ=0.1, 25° C.).

FIG. 136 shows the effect of frequency on normalized charges of Sm, Co and H2 in Co—Sm alloys electrodeposition (γ=0.1, 25° C.).

FIG. 137 shows XRD (left) and SEM (right) of deposits obtained from bath 1 at various frequencies (100 mA/cm2, γ=0.1, 25° C.).

FIG. 138 shows SEM of Co—Sm thin films obtained from bath 1 at 25° C., 2 kHz, γ=0.1 and at (a)-(b) 500 mA/cm2 and (c)-(d) 500 mA/cm2.

FIG. 139 shows the effects of frequency on magnetic properties at 25° C.

FIG. 140 shows the effect of Ton on (a) samarium deposit content and (b) current efficiency (period=100 ms, 25° C.).

FIG. 141 shows the effect of Ton on normalized charges of Sm, Co and H2 in Co—Sm alloys electrodeposition (period=100 ms, 25° C.).

FIG. 142 shows XRD (left) and SEM (right) of deposits obtained from bath 1 at various Ton (1000 mA/cm2, period=1000 ms, 25° C.).

FIG. 143 shows the effects of Ton on magnetic properties.

FIG. 144 shows Deposition rates of Sm and Co vs. deposition time in the electrodeposition of Co—Sm alloys.

Additional advantages of the invention will be set forth in part in the description which follows, and in part will be obvious from the description, or can be learned by practice of the invention. The advantages of the invention will be realized and attained by means of the elements and combinations particularly pointed out in the appended claims. It is to be understood that both the foregoing general description and the following detailed description are exemplary and explanatory only and are not restrictive of the invention, as claimed.

DETAILED DESCRIPTION

The present invention can be understood more readily by reference to the following detailed description of the invention and the Examples included therein.

Before the present compounds, compositions, articles, systems, devices, and/or methods are disclosed and described, it is to be understood that they are not limited to specific synthetic methods unless otherwise specified, or to particular reagents unless otherwise specified, as such may, of course, vary. It is also to be understood that the terminology used herein is for the purpose of describing particular aspects only and is not intended to be limiting. Although any methods and materials similar or equivalent to those described herein can be used in the practice or testing of the present invention, example methods and materials are now described.

All publications mentioned herein are incorporated herein by reference to disclose and describe the methods and/or materials in connection with which the publications are cited. The publications discussed herein are provided solely for their disclosure prior to the filing date of the present application. Nothing herein is to be construed as an admission that the present invention is not entitled to antedate such publication by virtue of prior invention. Further, the dates of publication provided herein can be different from the actual publication dates, which may need to be independently confirmed.

A. Definitions

As used in the specification and the appended claims, the singular forms “a,” “an” and “the” include plural referents unless the context clearly dictates otherwise. Thus, for example, reference to “a substrate,” “an alloy,” or “a sample” includes mixtures of two or more such substrates, alloys, or samples, and the like.

Ranges can be expressed herein as from “about” one particular value, and/or to “about” another particular value. When such a range is expressed, another aspect includes from the one particular value and/or to the other particular value. Similarly, when values are expressed as approximations, by use of the antecedent “about,” it will be understood that the particular value forms another aspect. It will be further understood that the endpoints of each of the ranges are significant both in relation to the other endpoint, and independently of the other endpoint. It is also understood that there are a number of values disclosed herein, and that each value is also herein disclosed as “about” that particular value in addition to the value itself. For example, if the value “10” is disclosed, then “about 10” is also disclosed. It is also understood that each unit between two particular units are also disclosed. For example, if 10 and 15 are disclosed, then 11, 12, 13, and 14 are also disclosed.

As used herein, the terms “optional” or “optionally” means that the subsequently described event or circumstance may or may not occur, and that the description includes instances where the event or circumstance occurs and instances where it does not.

Disclosed are the components to be used to prepare the compositions of the invention as well as the compositions themselves to be used within the methods disclosed herein. These and other materials are disclosed herein, and it is understood that when combinations, subsets, interactions, groups, etc., of these materials are disclosed that while specific reference of each various individual and collective combinations and permutation of these compounds may not be explicitly disclosed, each is specifically contemplated and described herein. For example, if a particular compound is disclosed and discussed and a number of modifications that can be made to a number of molecules including the compounds are discussed, specifically contemplated is each and every combination and permutation of the compound and the modifications that are possible unless specifically indicated to the contrary. Thus, if a class of molecules A, B, and C are disclosed as well as a class of molecules D, E, and F and an example of a combination molecule, A-D is disclosed, then even if each is not individually recited each is individually and collectively contemplated meaning combinations, A-E, A-F, B-D, B-E, B-F, C-D, C-E, and C-F are considered disclosed. Likewise, any subset or combination of these is also disclosed. Thus, for example, the sub-group of A-E, B-F, and C-E would be considered disclosed. This concept applies to all aspects of this application including, but not limited to, steps in methods of making and using the compositions of the invention. Thus, if there are a variety of additional steps that can be performed it is understood that each of these additional steps can be performed with any specific aspect or combination of aspects of the methods of the invention.

It is understood that the compositions disclosed herein have certain functions. Disclosed herein are certain structural requirements for performing the disclosed functions, and it is understood that there are a variety of structures that can perform the same function that are related to the disclosed structures, and that these structures will typically achieve the same result.

B. Rare Earth-Transition Metal (RE-TM) Permanent Magnets

RE-TM permanent magnets were developed at 1960's and became a practical commercial product at 1970. At that time, the RE-TM permanent magnets offered ten times higher coercivity and five times greater energy density than the best magnets of the 1960's. The maximum energy product for hard permanent magnetic materials is shown in FIG. 1 [G. J. Long and F. Grandjean, Supermagnets, hard magnetic materials, Kluwer Academic Publishers, Norwell, Mass., (1991), p3.].

RE-TM permanent magnets of different compositions exhibit a wide range of magnetic properties and cost; the research and development of Sm—Co and Nd—Fe-B alloys got more attention among these magnets for their superior performance and practical applications [K. J. Stmat, IEEE Trans. Magnetics, Mag-23, 2094, (1987).]. The first practical RE-TM permanent magnet, sintered SmCo5, was available about 1970 in the U.S. [M. G. Benz and D. L. Martin, Appl. Phys. Letters, 17, 176, (1970).]. Right after SmCo5, the investigation of quasi-binary intermetallics, RE2(Co, Fe)17 [A. E. Ray and K. J. Strnat, IEEE Trans. Magnetics, Mag-8, 516, (1972); K. J. Strnat, IEEE Trans. Magnetics, Mag-8, 511, (1972).], led to the development of the second generation of REIG permanent magnet —Sm2Co17. The first useful Sin2Co17 was developed in Japan in 1975 [T. Ojima, S. Tomizawa, T. Yoneyama, and T. Hori, Japan. J. Appl. Phys., 16, 671, (1977).]. The third generation of RE-TM permanent magnets, Nd2Fel4B, was developed by US and Japanese researchers and announced in 1983 [7; J. J. Croat, J. F. Herbst, R. W. Lee and F. E. Pinkerton, J. Appl. Phys., 55, 2078, (1984); N. C. Koon and B. N. Das, J. Appl. Phys., 55, 2063, (1984).] and brought new, much higher energy-product permanent magnets and a promise for cheaper RE-TM permanent magnets. The development of RE-TM permanent magnets not only resulted in a breakthrough of high performance magnetic materials but also created new applications for ferromagnetism materials.

Driven by the increasing interest of high performance permanent magnets, the improvement of RE-TM magnets has been mainly done by the incorporation and substitution of specific elements to RE-TM alloys. For example, small substitution of Cu for Co in SmCo5 leads to the precipitation of a nonmagnetic phase which increased the coercivity [E. A. Nesbitt, R. H. Willens, R. C. Sherwood, E. Buehler, and J. H. Wernick, Appl. Phys. Letters, 12, 361, (1968).]; the replacement of parts of Co by Fe in Sm2Co17 resulted in greater magnetization saturation (Ms) [A. E. Ray and K. J. Strnat, IEEE Trans. Magnetics, Mag-8, 516, (1972).] the replacement of 50% of Fe by Co in Nd2Fe14B gives higher Curie temperature [R. Grossinger, R. Krewenka, H. Buchner, and H. Harada, J. Phys. (Paris), 49, C8-659, (1988).]. These improvements make RE-TM permanent magnets easier to use for different kinds of industrial applications (Table 1) and enhances the performance of these magnetic devices.

Except for the traditional ways of making RE-TM magnets (i.e., bonding and sintering [M. G. Benz and D. L. Martin, Appl. Phys. Letters, 17, 176, (1970); M. G. Benz and D. L. Martin, J. Appl. Phys., 43, 4733, (1972); A. E. Ray and K. J. Strnat, IEEE Trans. Magnetics, Mag-11, 1429, (1975).]), new manufacture methods (i.e. mechanical alloying [J. Wecker, M. Katter and L. Schultz, J. Appl. Phys., 69, 6085, (1991); J. Ding, P. G. McCormick and R. Street, J. Alloys Comp., 191, 197, (1993).], nano powder metallurgy [J. Ding, Y. Liu, P. G. McCormick and R. Street, J. Magn. Magn. Mater., 123, L239, (1993).], and thin film processes -DC sputtering [H. C. Theuerer, E. A. Nesbitt, and D. D. Bacon, J. Appl. Phys., 40, 2994, (1969); S. A. Bendson and J. H. Judy, IEEE Trans. Magnetics, 9, 627, (1973); C. Zhang, R. Liu and G. Feng, IEEE Trans. Magnetics, 16, 1215, (1980); H. S. Cho, J. R. Salem, A. J. Kellock and R. B. Beyers, IEEE Trans. Magnetics, 33, 2890, (1997); R. Andreescu and M. J. O'Shea, I Appl. Phys., 91, 8183, (2002); 22], Rf sputtering [V. Neua and S. A. Shaheen, J. Appl. Phys., 53, 2401, (1982); F. J. Cadieu, S. H. Aly and T. D. Cheung, J. Appl. Phys., 64, 5501, (1988); K. Chen, H. Hegde and F. J. Cadieu, Appl. Phys. Letters, 61, 1861, (1992); T. Numata, H. Kinyama and S. Inokuchi, Appl. Phys., 86, 7006, (1999).], PVD [V. Geiss, E. Kneller and A. Nest, Appl. Phys., A27, 79, (1982); M. Gronau, H. Goeke, D. Schaffler and S. Sprenger, IEEE Trans. Magnetics, Mag-19, 1653, (1983); U. Kullmann, E. Koester and C. Dorsch, IEEE Trans. Magnetics, Mag-20, 420, (1984).], and pulsed laser deposition [V. Neu, J. Thomas, S. Faller, B. Holzapfel and L. Schultz, J. Magn. Magn. Mater., 242-245, 1290, (2002); F. J. Cadieu, R. Rani, and T. Theodoropoulos and Li Chen, J. Appl. Phys., 85, 5895, (1999).]) have been studied and developed. The improvement of manufacturing process not only promotes these high performance devices for traditional applications (i.e., automotive, domestic, electronic, aerospace devices) but also for new industrial applications, such as information storage, micro-electromechanical systems (MEMS), and nano-electromechanical systems (NEMS) of thin film RE-TM magnets.

TABLE 1
Applications of permanent magnets
ApplicationMagnetic Devices and Products
Automotivedc motor drivers, starter motors, window winders, wipers, fans,
speed meters, alternators
Domesticanalogues, watches, video recorders, electric clocks, hearing aids,
loudspeakers
Electronic andsensors, contactless switches, nmr spectrometers, energy meter
Instrumentationbearing, transducers, computer printer head, damper
Aerospacefrictionless bearings, couplings, magnetrons, klystrons, auto
compasses
Biosurgicaldentures, magnetic sphincters, magnetic sutures, cancer cell
separators, artifical hearts
Information storagemagneto-optical recording medium, perpendicular recording
media,
MEMS & NEMSmicromotor, actuator, magnetometer, magnetic sensors,
magnetic bubble memory

Recent developments of RE-TM magnets have focused on Co—Sm thin films by sputtering on the substrate with Cr underlayer [C. Prados and G. C. Hadjipanayis, J Appl. Phys., 83, 6253, (1998); C. Prados, A. Hernando, G. C. Hadjipanayis and J. M. Gonza'leza, J. Appl. Phys., 85, 6148, (1999); C. Prados and G. C. Hadjipanayis, Appl. Phys. Letters, 74, 430, (1999).]. Proper deposition conditions, alloy composition, and heat treatments increase the coercivity up to 40 kOe which is much higher than the coercivity of conventional SmCo5 (about 10 kOe) by other processes. In addition, using Cu as under layer in the sputtering process changes Co—Sm alloys from an in-plane to a perpendicular magnetic anisotropy [J Sayama, T. Asahi, K. Mizutani and T. Osaka, J. Phys. D: Appl. Phys., 37, L1, (2004); J. Sayama, K. Mizutani, T. Asahi, J. Ariake, K. Ouchi, S. Matsunuma and T. Osaka, J. Magn. Magn. Mater., 287, 239, (2005).], which makes it a good candidate for high density perpendicular recording media (also in terms of its excellent thermal stability and small minimal stable grain size).

A disadvantage of RE-TM permanent magnets to complete in the world market and wide use is their price [G. J. Long and F. Grandjean, Supermagnets, hard magnetic materials, Kluwer Academic Publishers, Norwell, Mass., (1991), pp 585-616.] which is strongly dependent on the manufacture process. Thin film processes, such as sputtering, PVD, and pulsed laser deposition, require a vacuum system and a high purity target to avoid impurity in deposits. In addition, the growth rates of these processes are slow. Therefore, making Co—Sm thin films by these processes is quite expensive and cannot provide any advantage in cost reduction making it difficult for commercialization. In other words, a cost effective manufacturing method must be developed to reduce the fabrication cost.

Electrodeposition is a simple, versatile and easily-controlled thin/thick film manufacturing method because of its simple setup, easy maintenance, low temperature operation, and low energy consumption. Compared to sputter, PVD, and other thin film processes, the most important advantage of electrodeposition is low cost. It was indicated that PVD process may be as much as ten times more expensive than electrodeposition [J. W. Dini, Plat. Surf: Finish., 80, 26, (1993).]. In addition, the growth rate of electrodeposition (5-0.1 μm/min) is a lot faster than other thin film processes (0.1-0.001 μm/min). This gives electrodeposition an advantage over other “thin film” technologies in thick film deposition which is often required in MEMS devices. With the help of masking patterns formed on the seedlayer, deposits of complex shape and geometry can be obtained by the electrodeposition. Therefore, electrodeposition is especially suitable to achieve high aspect ratio devices and microstructures in LIGA process [A. E. Ray and K. J. Strnat, IEEE Trans. Magnetics, Mag-8, 516, (1972).]. This provides electrodeposition more versatility and variety making it capable to adapt to various kind of applications. Therefore, using electrodeposition should effectively reduce the fabrication cost of RE-TM thin films and make them more competitive compared to other thin film technologies RE metal and alloys have been electrodeposited from molten salts [T. Iida T. Nohira and Y. Ito, Electrochim. Acta, 48, 901, (2003); T. Iida T. Nohira and Y. Ito, Electrochim. Acta, 48, 901, (2003).[P. Liu, Y. Du, Q. Yang, Y. Tong and G. A. Hope, J. Magn. Magn. Mater., 153, C57, (2006).] and nonaqueous solutions [Y. Sato, H. Ishida, K. Kobayakawa and Y. Abe, Chem. Lett., (8), 1471, (1990); Y. Sato, T. Takazawa, M. Takahashi, H. Ishida and K. Kobayakawa, Plat. Surf Finish., 80, 72, (1993).]. Unfortunately, most of the deposits obtained from non-aqueous media have poor magnetic properties, as low coercivity and saturation magnetization; oxides and hydroxides also be found in some cases after heat treatments [Y. Sato, T. Takazawa, M. Takahashi, H. Ishida and K. Kobayakawa, Plat. Surf Finish., 80, 72, (1993).]. On the other hand, few studies of the electrodeposition of IG-RE alloys from aqueous solutions have been reported [L. Chen, M. Schwartz, and K. Nobe, in Electrodeposited Thin Films, M. Paunovic and D. A. Scherson, Editors, PV 96-19, p. 239, The Electrochemical Society Proceedings Series, Pennington, N.J. (1996); Schwartz et al., in Magnetic Materials, Processes, and Devices V. Applications to Storage and Microelectromechanical Systems (MEMS), Romankiw et al., Editors, PV 98-20, p. 646, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999); M. Schwartz, N. V. Myung, and K. Nobe, J. Electrochem. Soc., 151, C468, (2004).].

Electrodeposition of RE metals from aqueous solutions is more difficult compared to non-aqueous solutions, as a result of anticipated vigorous hydrogen evolution at the reduction potentials of RE metals in aqueous solution. The reduction potentials of RE metals are extremely negative (E°<−2VsitE) [W. M. Latimer, The Oxidation States of the Elements and Their Potentials in Aqueous Solution, Prentice-Hall, New York, pp. 286-295 (1952).],muchlowerthanthereductionpotentialofwater (H2O+2e20H+1/2H2, E-0.826V). Therefore, instead of RE metal deposition, water can decompose first. In addition, RE metal ions would be hydrolyzed at PH>6 and tends to react with dissolved oxygen or hydroxyl ions to form oxide or hydroxide. Therefore, hydroxides and oxides would be deposited instead of RE metal making the deposition of RE metal from aqueous solution difficult, similar to electrodeposition of Mo, W and V from aqueous solutions. However, Mo, W and V has been co-electrodeposited with iron group metals from aqueous solution the addition [L. O. Case and A. Krohn, J. Electrochem. Soc., 105, 512 (1958); V. B. Singh, L. C. Singh, and P. K. Tikoo, J. Electrochem. Soc., 127, 590 (1980); M. Schwartz, Unpublished data, 1946; see also discussion in Trans. Electrochem. Soc., 94, 382 (1948); A. Brenner, P. Burkhead, and E. Seegmiller, J. Res. Natl. Bur. Stand., 93, 351 (1947); M. L. Holt and L. E. Vaaler, Trans. Electrochem. Soc., 94, 50 (1948);W. E. Clark and M. L. Holt, Trans. Electrochem. Soc., 94, 244 (1948); M. H. Lietzke and M. L. Holt, Trans. Electrochem. Soc., 94, 252 (1948); W. H. Safranek and L. E. Vaaler, Plating (East Orange, N.J.), 46, 133 (1959); Arcos et al., Magnetic Materials, Processes, and Devices IV. Applications to Storage and Microelectromechanical Systems (MEMS), Romankiw and Herman, Jr., Editors, PV 95-18, p. 563, The Electrochemical Proceedings Series, Pennington, N.J. (1996); Arcos et al., Plat. Surf Finish., 90 46 (2003).] of an appropriate complexer. Various hypotheses has been reviewed by Brenner [A. Brenner, Electrodeposition of Alloys, Vol. 2, pp. 400-453, Academic Press, NewYork (1963).] to explain the phenomenon of the co-reduction of these metals, which he referred to as “induced” co-deposition.

In 1947-1948, Schwartz initiated a commercial installation of a Co-W ammonium citrate plating process [J Sayama, T. Asahi, K. Mizutani and T. Osaka, J. Phys. D: Appl. Phys., 37, L1, (2004).]. He found that when solutions of Co and W salt were mixed, cobalt tungstate precipitate immediately but dissolved with the addition of citrate. He conjectured that both Co2+ and W6+ are present in the same complex with deportonation of the hydroxycarboxylate portions, resulting in a heteronuclear biscitrate complex.

In 1994, Schwartz initial research of the electrodeposition of IG-RE alloys from aqueous solution [M. Schwartz, Unpublished data, UCLA 1994] and tried to extend his idea of complex formation in aqueous Co—W alloy electrodeposition to IG-RE alloys. Between 1996 and 2004, a series of studies of IG-RE alloys from aqueous solutions were reported [L. Chen, M. Schwartz, and K. Nobe, in Electrodeposited Thin Films, M. Paunovic and D. A. Scherson, Editors, PV 96-19, p. 239, The Electrochemical Society Proceedings Series, Pennington, N.J. (1996); Schwartz et al., in Magnetic Materials, Processes, and Devices V. Applications to Storage and Microelectromechanical Systems (MEMS), Romankiw et al., Editors, PV 98-20, p. 646, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999); M. Schwartz, N. V. Myung, and K. Nobe, J. Electrochem. Soc., 151, C468, (2004); Myung et al., in Fundamental Aspects of Electrochemical Deposition and Dissolution, Matlosz et al., Editors, PV 99-33, p. 263, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999).]. The experimental results indicated that RE and IG metals can be co-deposited by addition of appropriate complexes, such as glycine and its derivatives, into the solution. Initially, studies were mainly focused on RE mixtures [L. Chen, M. Schwartz, and K. Nobe, in Electrodeposited Thin Films, M. Paunovic and D. A. Scherson, Editors, PV 96-19, p. 239, The Electrochemical Society Proceedings Series, Pennington, N.J. (1996).] for co-deposition of IG-RE alloys [Schwartz et al., in Magnetic Materials, Processes, and Devices V. Applications to Storage and Microelectromechanical Systems (MEMS), Romankiw et al., Editors, PV 98-20, p. 646, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999); M. Schwartz, N. V. Myung, and K. Nobe, J. Electrochem. Soc., 151, C468, (2004).]. The research focus of this dissertation is on the co-deposition of Co—Sm alloys to fabricate high performance SmCo5 and Sm2Co17 permanent magnets, which promise much lower costs, more flexibility than existing manufacturing processes.

C. Cobalt-Samarium Magnets

Co—Sm magnets are known for magnetocrystalline anisotropy, high coercivity, and maximum energy products. These characteristics are based mainly on the intermetallic phases, SmCo5 and Sm2Co16. SmCo5 magnets have the highest uniaxial anisotropies of any class of magnets, Ku (uniaxial magnetic anisotropy energy coefficient)≅107 J/m3. On the other hand, Sm2Co17 magnets exhibit high flux density and Curie temperature. Metastable phases, intermetallic compounds, and crystalline structures of Co—Sm alloys are the important properties that provide superior permanent magnet performance. The unique properties of SmCo5 and Sm2Co17 magnets will be reviewed.

Different kinds of manufacturing methods for fabricating Co—Sm magnets have been developed the past 40 years. Bonding and sintering [M. G. Benz and D. L. Martin, Appl. Phys. Letters, 17, 176, (1970); M. G. Benz and D. L. Martin, J. Appl. Phys., 43, 4733, (1972); A. E. Ray and K. J. Strnat, IEEE Trans. Magnetics, Mag-11, 1429, (1975); Z. A. Abdelnour, H. F. Mildrum and K. J. Strnat, IEEE Trans. Magnetics, Mag-16, 1980, (1980)] and mechanical alloying [J. Wecker, M. Katter and L. Schultz, J. Appl. Phys., 69, 6085, (1991); S. K. Chen, J. L. Tsai and T. S. Chin, J. Appl. Phys., 81, 5631, (1997); M. L. Kahn, J. L. Bobet, F. Weil and B. Chevalier, J. Alloys Comp., 334, 285, (2002); J. Zhou, R. Skomski and D. J. Sellmyer, J. Appl. Phys., 93, 6495, (2003)] are the main methods for processing large Co—Sm magnets for common applications (e.g., automotive, domestic, electronic, aerospace devices). Applications for information storage, microelectromechanical systems (MEMS), and nanoelectromechanical systems (NEMS) have lead to developments of thin film processes such as DC sputtering, RF sputtering, PVD, and pulsed laser deposition. The preparation of thin film Co—Sm magnets not only lead to new industrial applications but also improve performance. Thin film processes for Co—Sm magnets are disclosed.

1. Co—Sm Alloy System

The potential superior magnetic properties of intermetallic compounds of Sm and Co was the focus of Buschow et al.'s [K. H. J. Buschow and W. A. J. J. Velge, J. Less-Common Met., 13, 11, (1967); K. H. Buschow and A. S. V. D. Goot, J. Less-Common Met., 14, 323, (1967)] investigation of the entire concentration range of binary Sm—Co alloys by X-ray diffraction, thermoanalytical and metallographic methods. The phase diagram is shown in FIG. 2. Intermetallic compounds of Sm3Co, Sm9Co4, SmCo2, SmCo3, Sm2Co7, SmCo5 and Sm2Co17 were obtained.

Three eutectics between the phases Sm—Sm3Co, Sm9Co4—SmCo2 and Sm2Co17—Co were found. FIG. 3 shows the Co-rich region of the Co—Sm phase diagram [K. J. Strnat, Ferromagnetic materials, Vol 4, E. P. Wohlfarth and K. H. J. Buschow, Editors, Elsevier Science Publishers, New York, (1998), p. 143] after den Broeder and Buschow (1972), Perry (1977), and Ray (1986). From the phase diagram, both SmCo5 and Sm2Co17 are identified as intermetallic compounds. Sm2Co17 is a stable phase and its melting point is about 1335° C. On the other hand, SmCo5 is a metastable phase existing above 805° C. (eutectoid point).

2. Properties of SMCO5 and SM2CO17 Magnets

The hexagonal CaCu5 structure of SmCo5 [C. Barrett and T. B. Massalski, Structure of metals, Pergamon Press, Oxford, New York, (1980), p 266] is shown in FIG. 4. A smaller hexagonal Co combined within a hexagonal Sm in the same plane (layer A and C). A similar hexagonal Co plane rotated by 30° (layer B) is inserted between two hexagonal Sm planes (layer A and C).

Single-phase SmCo5 [R. C. O'Handley, Modern magnetic materials, John Wiley & Son, Inc., New York, (2000), pp. 496-502] has exhibited room temperature coercivity (Hc) of 0.72 MA/m (9 kOe), maximum energy product of over 200 kJ/m3 (24 MG·Oe) and saturation magnetization (Ms) of 1 T. A high Curie point (Tc=685° C.) enables a wide range of applications. The high coercivity of SmCo5 is attributed to reversal domain nucleation controlled by grain boundaries [J. D. Livingston, AIP Conf Proc., 10, 643 (1973)] restricting the mobility of domain walls leading to high coercivity.

The rhombohedral Th2Ni17-type structure of Sm2Co17 [R. C. O'Handley, Modern magnetic materials, John Wiley & Son, Inc., New York, (2000), pp. 496-502] is shown in FIG. 5. This structure has the same cobalt hexagonal nets as SmCo5 but with fewer Sm atoms in adjacent layers. The magnetocrystalline anisotropy of Sm2Co17 magnets (Ku≅3×106 J/m3) is less than SmCo5 magnets (K≅107 J/m3) as well as its coercivity of 0.68 MA/m (8.5 kOe). Saturation magnetization of 1.2-1.5 T (109-137 emu/g) and Curie temperature (Tc) of 810-970° C. are higher, however. The grains of Sm2Co17 magnets containing fine structures of SmCo5 pin the domain walls [P. Campbell, Permanent magnet materials and their application, Cambridge University Press, New York, (1994), pp. 42-43] and can be achieved by proper heat treatment. The pinning mechanism is shown as FIG. 6.

Chemical reduction is usually used to produce Co—Sm powders [P. Campbell, Permanent magnet materials and their application, Cambridge University Press, New York, (1994), pp. 38-45]. Co—Sm magnets can be obtained by sintering Co—Sm powders, molding bonded Co—Sm particles, mechanical alloying, and thin film processes, such as sputter and PVD. After optimization of heat treatment, high coercivities and anisotropic magnetic properties can be achieved.

3. Thin Film Technology of Co-Sm Magnets

a. DC Sputtering

The first Co—Sm thin film was prepared in 1969 by getter sputtering [H. C. Theuerer, E. A. Nesbitt, and D. D. Bacon, J. Appl. Phys., 40, 2994, (1969)] for comparison with the unusually high coercivity (28.7 kOe, annealed at 400° C.) of sintered Co5-xCuxSm alloys found by E. A. Nesbitt et al. [E. A. Nesbitt, R. H. Willens, R. C. Sherwood, E. Buehler, and J. H. Wernick, Appl. Phys. Letters, 12, 361, (1968).]. The later obtained Co—Sm thin films, with/without addition of Cu, with very high coercivities were dependent strongly on substrate temperatures (bell-shaped curves), and higher than bulk specimens.

Theuerer et al. reported a maximum coercivity of 20 kOe for SmCo5 obtained with films (400 nm) prepared at 600° C. which is much higher than the value of only 1 kOe for bulk specimens of the same composition [H. C. Theuerer, E. A. Nesbitt, and D. D. Bacon, J. Appl. Phys., 40, 2994, (1969)]. For the Co5-xCuxSm (x=1.35) alloys, films (400 nm) prepared at 500° C. had a maximum coercivity of 30 kOe which is greater than bulk specimens (12 kOe). Thick films (˜5 um) prepared at 500-600° C. for the respective alloy compositions had lower coercivities (13.3 kOe) than thin films (30 KOe, 400 nm) and close to the bulk values (12 KOe). Films deposited on crystalline substrates resulted in larger grain size with lower coercivities.

Bendson and Judy [S. A. Bendson and J. H. Judy, IEEE Trans. Magnetics, 9, 627, (1973)] used DC triode sputtering to obtain Co—Sm thin films (100 to 500 nm) at 10−3 Torr of Ar on glaze alumina substrates at or above 600° C. from Co and Sm targets. Saturation magnitization, coercivity and squareness of deposits are shown in FIG. 7. The saturation magnitization is reduced by increasing Sm content until the mixture becomes paramagnetic between 25 and 30 atomic percent Sm (FIG. 7) in agreement with the published value for bulk specimens [R. Lemaire, R. Pauthenet and J. Schweizer, IEEE Trans. Magnetics, Mag-6, 153, (1970)]. In addition, low coercivity is observed over the major range of composition except among 22 at % Sm where extremely high coercivity is found. In these films, the easy axis lies in the film plane. Addition of Sm apparently results in decreased squareness.

Zhang et al. [C. Zhang, R. Liu and G. Feng, IEEE Trans. Magnetics, 16, 1215, (1980)] studied non-crystalline SmCo5 and Sm0.5MM′0.5Co (MM′: Ce 50%, La 20%, Nd 10%, Pr 10%) thin films (0.7 to 1.4 m) by DC diode sputtering on a liquid nitrogen cooled glass substrate. Even after heat treatment at 750° C. for 4 hours, crystalline SmCo5 was not formed. The coercivity is only about several hundred Oe which is far below 20 KOe obtained by Theuerer et al. [H. C. Theuerer, E. A. Nesbitt, and D. D. Bacon, J. Appl. Phys., 40, 2994, (1969)]. On the other hand, due to the formation of crystalline (β-Co (fcc), saturation magnetization of SmCo5 non-crystalline films were about 620-663G (74-79 emu/g) and increased to 810-880 G (97-105 emu/g) after heat treatment at 600° C. for 2 hours.

Co—Sm thin films (79 at % Co, 21 at % Sm, 24 nm) on Cr (95 nm) were first studied in 1994 by DC magnetron sputtering (no heat treatment) [Y. Liu, B. W. Robertson, Z. S. Shan, S. Malhotra, M. J. Yu, S. K. Renukunta, S. H. Liou and D. J. Sellmyer, IEEE Trans. Magnetics, 30, 4035, (1994)]. It was found that the volume fraction of the crystallites decreased from 91 to 54 at % as the Ar pressure increased from 0.5×10−4 to 3×10−4 Torr the maximum coercivity of 2.58 kOe was obtained at 1.2×10−4 Torr.

Mizukami et al. [M. Mizukami, T. Abe and T. Nishihara, IEEE Trans., Magnetics, 33, 2977, (1997)] found that the coercivity of DC sputtered Co—Sm thin film decreased by 50% after exposure to air for 30 hours because Sm was oxidized. A protective Cr layer coating reduced oxidation by forming a Cr/CoSm/Cr structure. Takei et al. [S. Takei, A. Morisako and M. Matsumoto, J. Appl. Phys., 81, 4674, (1997)] found that the coercivity of Co—Sm (20 at % Sm) thin films increased linearly with the increasing Ar pressure. The squareness ratio of 0.92 could be obtained with a highly crystallized Cr underlayer (PAr of 10−3 Torr). These values suggest that an easy axis of magnetization for the CoSm thin film is in-plane. The coercivity and squareness (1.95 kOe and 0.92, respectively) of CoSm thin film (20 at % Sm) with Cr underlayer in this study is higher than that obtained by Bendson and Judy (0.7 kOe and 0.6) [S. A. Bendson and J. H. Judy, IEEE Trans. Magnetics, 9, 627, (1973)].

The effects of oxidation on the magnetic and electrical properties of DC sputtered CoSm thin films were studied by Cho et al. [H. S. Cho, J.R. Salem, A.J. Kellock and R. B. Beyers, IEEE Trans. Magnetics, 33, 2890, (1997)]. Co—Sm films prepared on Si (100) and quartz glass by DC magnetron sputtering of different composition and substrate temperatures were characterized. The results are summarized in FIG. 8. Co—Sm films deposited at room temperature were non-crystalline, and the coercivity increased from 20 to 300 Oe and the magnetization saturation decreased from 1400 to 500 emu/cm3 (154 to 60 emu/g or 1.76 to 0.63 T) with increasing Sm deposit content from 0 to 28 at %. With increased substrate temperature from 25 to 350° C., the coercivity increased from 200 to 8000 e and the magnetization saturation decreased from 500 to 380 emu/cm3 (60 to 46 emu/g or 0.63 to 0.48 T).

Cho et al. believed that these different behaviors depend on the extent of oxidation in the film especially at high substrate temperatures [H. S. Cho, J. R. Salem, A. J. Kellock and R. B. Beyers, IEEE Trans. Magnetics, 33, 2890, (1997)]. More than 10 at % oxygen was present in sputtered deposits at room temperature, and the amount of oxygen increased with substrate temperature. TEM results indicated that Co—Sm thin films deposited at 360° C. were multiphase mixtures of Sm2O3 and Co-enriched phases rather than a simple homogeneous Co—Sm phase.

In 1998 Liu et al. [Y. Liu, R. A. Thomas, S. S. Malhotra, Z. S. Shan, S. H. Liou and D. J. Sellmyer, J. Appl. Phys., 83, 6244, (1998)] studied phase formation and magnetic properties of Co—Sm thin films (DC magnetron sputtering) by increasing thickness with a Cr cover layer as Mizukami et al [M. Mizukami, T. Abe and T. Nishihara, IEEE Trans., Magnetics, 33, 2977, (1997)] compared to his previous work [Y. Liu, B. W. Robertson, Z. S. Shan, S. Malhotra, M. J. Yu, S. K. Renukunta, S. H. Liou and D. J. Sellmyer, IEEE Trans. Magnetics, 30, 4035, (1994)]. It was found that for deposits (19 at % Sin, 360 nm) obtained at Ar pressure of 2×10−4 Torr then annealed at 600° C. has SmCo5 phases and an extremely high coercivity of 45 kOe. TEM showed a new phase SmCo3 (22 at % Sm) was formed when the film was annealed at 500° C.

Prados et al [C. Prados and G. C. Hadjipanayis, J. Appl. Phys., 83, 6253, (1998); C. Prados and G. C. Hadjipanayis, Appl. Phys. Letters, 74, 430, (1999); C. Prados, A. Hernando, G. C. Hadjipanayis and J. M. Gonzaleza, J. Appl. Phys., 85, 6148, (1999)] sputtered Sm(Co, Ni, Cu) thin films (500 nm) on Cr underlayers (300 nm) on water cooled Si to promote a c-axis texture along the in-plane direction in order to increase their in-plane magnetic anisotropy. After films obtained from DC magnetron sputtering, the thin films were annealed in a vacuum greater than 10−5 Torr at 400 to 650° C. for 30 min. Crystallization of the non-crystalline magnetic films (obtained at room temperature) produced a huge enhancement of coercivity (from 100 to more than 40 kOe). High angle XRD patterns indicate that the Cr underlayers grew with a (110) texture, and the Sm(Co, Ni, Cu) thin films were non-crystalline as deposited. After annealing, the (111) plane of SmCo5 appeared and optimum samples indicated a nanocrystalline structure (particle size of 10 nm).

After Liu and Prados showed that the Cr underlayer plays a significant role in the magnetic properties of Co—Sm thin films, Takei et al. [S. Takei, A. Morisako and M. Matsumoto, J. Appl. Phys., 87, 6968, (2000)] worked on the effect of different kinds of underlayers, such as Cr, Mo, W, W/Cr and Al. The coercivity of the films with Cr and Mo underlayer were larger than 3 kOe for underlayers thicker than 100 nm. The squareness of the films with Cr and Mo underlayers were higher than 0.85 and 0.94, respectively. The result indicated that both Cr and Mo are suitable for SmCo films with in-plane magnetic anisotropy. For ultra-thin Co—Sm films (2.5 nm) deposited on Cr underlayer (100 nm) on Corning #7059 glass the XRD pattern indicated that the Sm—Co layer crystallized by substrate heating with coercivity higher than 3 kOe. Substrate heating during deposition was effective in preparing ultrathin Sm—Co thin films with higher coercivity.

There have been many in-plane magnetic anisotropy studies of Co—Sm thin films but few of perpendicular magnetic anisotropy. Recently, Sayama et al. [J. Sayama, T. Asahi, K. Mizutani and T. Osaka, J. Phys. D: Appl. Phys., 37, L1, (2004); J. Sayama, K. Mizutani, T. Asahi, J. Ariake, K. Ouchi, S. Matsunuma and T. Osaka, J. Magn. Magn. Mater., 287, 239, (2005)] DC sputtered CoSm thin films [Co(0.41 nm)/Sm(0.31 nm)]35 on Cu underlayer (100 nm) on a glass substrate at various temperatures. Coercivity increased rapidly between 325-345° C. and was greater in the perpendicular direction than the in-plane direction. The Co/Sm laminate structure on Cu underlayer was the key in the crystallization of SmCo5 with its c-axis perpendicular to the film plane. The perpendicular magnetic anisotropy was further improved by reducing the surface roughness of the Cu underlayer.

Takei et al. [S. Takei, A. Morisako and M. Matsumoto, J. Magn. Magn. Mater., 272-276, 1703, (2004)] obtained similar results to Sayama et al. [J. Sayama, T. Asahi, K. Mizutani and T. Osaka, J. Phys. D: Appl. Phys., 37, L1, (2004); J. Sayama, K. Mizutani, T. Asahi, J. Ariake, K. Ouchi, S. Matsunuma and T. Osaka, J. Magn. Magn. Mater., 287, 239, (2005)] and found that the Co—Sm layer with (001) orientation was crystallized on top of the (111) Cu orientation underlayer and perpendicular coercivity of the film was about 9.6 kOe with substrate temperatures about 300° C.

b. RF-Sputtering

Cadieu et al. [F. J. Cadieu, S. H. Aly and T. D. Cheung, J. Appl. Phys., 53, 2401, (1982)] deposited Co—Sm thin films (1.61 lam) deposited on an Al2O3 substrate by RF-sputtering. The composition of the Co—Sm thin films were close to the composition SmCo5 and Sm2Co17. Hysteresis loops for various Sm deposit content are shown in FIG. 9.

The low saturation magnitization measured perpendicular to the film plane and the XRD patterns indicated that the c-axis is strongly oriented in the film plane direction. This result is similar to the crystalline Co—Sm thin films made by DC sputtering onto heated substrate or by crystallization (annealing) of a noncrystalline deposit. The energy product measured in the in-plane direction was lower than bulk specimens which might be due to the random orientation of the c-axis in the film plane.

Velu et al. [E. M. T. Velu and D. N. Lambeth, J Magn. Magn. Mater., 69, 5175, (1991)] studied the Co—Sm thin films with/without Cr underlayer on 7059 Coming glass and NiP coated Al substrates. The coercivity of the CoSm thin films was higher for Cr underlayers deposited at 10−2 Torr Ar pressure. The coercivity decreased with higher substrate temperature (>300° C.) which might due to the crystallographic texture transformation of Cr from <110> to <200> orientation. A maximum coercivity of 2.4 kOe and squareness of 1 were obtained for CoSm thin films (14nm) under optimal conditions.

Chen et al. studied the induced anisotropy [K. Chen, H. Hegde and F. J. Cadieu, Appl. Phys. Letters, 61, 1861, (1992)] and different other types of anisotropy [K. Chen, H. Hegde, S. U. Jen and F. J. Cadieu, J Appl. Phys., 73, 5923, (1993)] in RF sputtered non-crystalline Sm—Co thin films on water cooled polycrystalline Al2O3 substrates. An in-plane magnetic field was applied during RF sputtering. Both in-plane and perpendicular anisotropy were found depending on the sputtering conditions. Three different sources of anisotropy can be distinguished in these films. The in-plane anisotropy was explained as directional pair ordering; perpendicular anisotropy was only observed for films deposited through sputtering at room temperature; a much larger anisotropy was observed at higher deposition temperatures with the easy axis in the film plane.

Neu et al. [V. Neua and S. A. Shaheen, J. Appl. Phys., 86, 7006, (1999)] RF-sputtered SmCo5 and Sm(CoFeCuZr)7 thin films (1 μm) on heated polycrystalline Al2O3 substrates. With increased Sm deposit content, Sm—Co thin films transformed from the TbCu7-type to the CaCu5-type structure at around 17 at % Sm in agreement with the stoichiometric SmCo5 (16.7 at %). With increasing Sm content, the c-axis of SmCo5 preferred to lie on the film plane. The morphology showed larger and more elongated grains with increasing Sm content which might be due to a higher surface mobility of Sm-rich samples. Higher Sm content resulted in higher coercivity and lower Mrperpendicular/Mr parallel values (FIG. 10), the result of crystallization of SmCo5.

c. PVD

Geiss et al. [V. Geiss, E. Kneller and A. Nest, Appl. Phys., A27, 79, (1982)] studied non-crystalline Sm100-xCox (70<x<90) thin films (150 nm) prepared by vapor deposition on flat glass substrates at room temperature. A magnetic field of 500 Oe was applied parallel to the film plane during evaporation. After deposition, specimens were aged at various temperatures. The easy axis of magnetization lay on film plane and deposits appeared non-crystalline magnets. The coercivity varied between 30 to 3000 Oe, depending on the composition, temperature and heat treatment. Aging of any sample at temperatures below the crystallization temperature resulted in a decrease in coercivity.

Gronau et al. [M. Gronau, H. Goeke, D. SchUffler and S. Sprenger, IEEE Trans. Magnetics, Mag-19, 1653, (1983)] prepared non-crystalline Sm1-xCox (0.67<x<0.91) thin films (10 to 350 nm) by flash-evaporation of SmCo-alloy powder on glass substrates. Saturation magnitization decreased linearly with increasing Sm content with the same slope as the crystalline material, but differences were about 10% smaller for the non-crystalline phase in agreement with the results of DC sputtered deposits by Bendson and Judy [S. A. Bendson and J. H. Judy, IEEE Trans. Magnetics, 9, 627, (1973)]. For x=0.67 (SmCo2) crystals were nonmagnetic with Ms nearly zero. The coercivity increased, reached a maximum of 53 kA/m (650 Oe) at x=0.74, then decreased. They concluded that there was little difference for films prepared by evaporation on heated substrate or annealed at the same temperature.

Following Geiss [K. Chen, H. Hegde, S. U. Jen and F. J. Cadieu, J Appl. Phys., 73, 5923, (1993)] and Gronau [M. Gronau, H. Goeke, D. SchUffler and S. Sprenger, IEEE Trans. Magnetics, Mag-19, 1653, (1983)], Kullmann et al. [U. Kullmann, E. Koester and C. Dorsch, IEEE Trans. Magnetics, Mag-20, 420, (1984)] prepared non-crystalline Sm100-xCox (75<x<90) thin films (100 nm) by PVD to obtain high density longitudinal recording. With increasing Sm content, the coercivity increased from 30 to 100 kA/m (375 to 1250 Oe). The saturation magnetization was decreased with decreased Sm content.

The recording performance showed an improvement in recording density and signal-to-noise ratio compared to traditional longitudinal recording media.

4. Electrodeposition of Co-RE (Rare Earth) Alloys

a. Electrodeposition of Co-RE Alloys from Non-Aqueous Solutions

In 1953, Moeller et al. [T. Moeller and P. A. Zimmerman, J. Am. Chem. Soc, 75, 3940, (1953); T. Moeller and P. A. Zimmerman, Science, 120, 539, (1954)] dissolved anhydrous yttrium acetate, neodymium bromide and lanthanum nitrate in anhydrous ethylenediamine and monoethanolamine to obtain yttrium, neodymium and lanthanum. Electrolyses of ethylenediamine solutions gave metallic cathode deposits with all salts tested, but deposits were not obtained from monoethanolamine solutions due to low solubility and conductivities. All deposits exhibited rare-earth metal properties, such as oxidation in air or in water and hydrogen evolution from hydrochloric acid solution.

Increased interest in higher performance Co—Sm magnets (SmCo5 and Sm2Co17) lead to studies of the electrodeposition of Co-Sin alloys from non-aqueous media since aqueous electrodeposition of Co—Sm alloys was extremely difficult, if not impossible, due to excessive hydrogen evolution. [Y. Sato, H. Ishida, K. Kobayakawa and Y. Abe, Chemistry Letter, 1471 (1990); T. Lida, T. Nohira and Y. Ito, Electrochim. Acta, 48, 901, (2003); T. Lida, T. Nohira and Y. Ito, Electrochim. Acta, 48, 2517, (2003); P. Liu, Y. Du, Q. Yang, Y. Tong and G. A. Hope, J. Magn. Magn. Mater., 153, C57, (2006)]

Sato et al. [Y. Sato, T. Takazawa, M. Takahashi, H. Ishida and K. Kobayakawa, Plat. Surf Finish, 72, 80, (1993); Y. Sato, H. Ishida, K. Kobayakawa and Y. Abe, Chemistry Letter, 1471 (1990)] electrodeposited non-crystalline Sm—Co thin films from a formamide solution containing anhydrous samarium and cobalt chloride. The deposits were confirmed as metallic by XPS. The results suggested that the deposit contains Co-rich compounds such as Sm2Co17 and SmCo3, which show ferromagnetism. However, after deposits were heat treated at 600° C. for 3 hours, cobalt oxides were found in the specimens. The magnetic properties of deposits after annealing did not improve. The highest coercivity before heat treatment was about 90 Oe and after heat treatment at 600° C. for 3 hours about 562 Oe.

In 2003, Iida et al. studied the electrodeposition of Sm—Co alloys at a Co cathode in a molten LiCl—KCl—SmCl3 system [T. Lida, T. Nohira and Y. Ito, Electrochim. Acta, 48, 901, (2003)] at 723 K. In addition, a molten LiCl—KCl—SmCl3—CoCl2 system [T. Lida, T. Nohira and Y. Ito, Electrochim. Acta, 48, 2517, (2003)] at 450° C. using a Cu substrate also had been studied. Phases of the deposited Sm—Co alloys could be controlled by the potential. Sm2Co17, SmCo3, SmCo2, and LixSm4Co6 were found at the potentials of 1.4, 0.8, 0.3 and 0.05V (vs. Li+/Li). However, the deposit rate was slow (0.004-5 um/hr) and magnetic properties were not measured.

Recently, in 2006, Liu et al. [P. Liu, Y. Du, Q. Yang, Y. Tong and G. A. Hope, J. Magn. Magn. Mater., 153, C57, (2006)] codeposited Sm—Co non-crystalline films in a urea-acetamide-NaBr-MCl melt (M=Sm or Co). They found that the reduction of Co is irreversible and Sm cannot be reduced alone in these melts and applied the idea of a polynuclear complex mechanism as proposed by Schwartz et al [N. V. Myung, M. Schwartz, and K. Nobe, in Fundamental Aspects of Electrochemical Deposition and Dissolution, M. Matlosz, D. Landolt, R. Aogaki, Y. Sato, and J. B. Talbot, Editors, PV 99-33, p. 263, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999).] in the glycine aqueous system. The Sm—Co films were non-crystalline as deposited, and intermetallic Sm—Co phases were found after heat treatment at 900° C. Unfortunately, the saturation magnetization of a deposit of 3.3 at % Sm is only 2.9 emu/g (before heat treatment) which is too low compared to DC sputtering (143 emu/g) for a similar alloy composition. The highest coercivity of deposits after heat treatment (900° C., 3 hr) was 180 Oe (60 at % Sm) with saturation magnetization of 1.66 emu/g.

b. Electrodeposition of IG-W, Mo and V Alloys

Brenner suggested that although metals with extreme negative reduction potentials, such as W, Mo and V, cannot be deposited from aqueous media individually, co-deposition with the iron group metals can be achieved. Review of representative investigations of these alloys follow:

c. IG-W Alloys

Holt and his student [J. Seim and M. L. Holt, Trans. Electrochem. Soc., 95, 205 (1949); D. W. Ernst, R. F. Amlie and M. L. Holt, J. Electrochem. Soc., 102 (8), 461 (1955); D. W. Ernst and M. L. Holt, Ibid., 105 (11), 686 (1958).] co-electrodeposited IG-W alloys from ammoniacal solutions containing organic hydroxyl acids, such as citric, tartric, and malic acid. Hydroxyl acids, which are known as good complexing agents, increased current efficiency and tungstate solubility to obtain smooth IG-W deposits. The reduction of the tungstate was suggested due to the catalysis of atomic hydrogen reduction by a two-step reduction hypothesis.

Brenner et al. [A. Brenner, P. Burkheard and E. Seegmiller, J. Res. NBS, 94, 351 (1947).] developed a codeposition process of W and iron group alloys from aqueous ammoniacal solutions (˜pH8.5) containing appropriate metal salts (sodium tungstate and 1G (iron group)-chloride or IG-sulfate) and certain hydroxyl-organic acids (citric acid, tartaric acid, hydroxyacetic acid, malic acid, gluconic acid). They concluded that the W ion concentration is the most important variable affecting W deposit content, which increased and reached a limit with increasing W concentrations. The maximum W deposit content of Co—W and Fe—W alloys was about 23 to 32 at %, and about 12 at % for Ni—W alloys. W deposit content increased with CD, but did not significantly change with temperature. However, current efficiency increased considerably with increase in temperature.

Schwartz (1948) developed a commercial Co—W plating process using an ammoniacal citrate bath [M. Schwartz, unpublished data, 1948-55; see also discussion in Trans. Electrochem. Soc., 94, 382 (1948)]. He found that when the Co and. WO42− salt solutions are mixed, a cobalt tungstate precipitate forms which dissolves with addition of citrate. His experimental results lead him to conjecture that both Co(I1) and W(VI) are coordinated in the same complex with deprotonation of the carboxylate forming a heteronuclearbiscitrato complex [M. Schwartz and K. Nobe, Trans. Electrochem. Soc., 1, 103 (2006)].

Recently, Gileadi and his coworkers [Younes and E. Gileadi, J. Electrochem. Soc., 149, C100, (2002); Younes-Metzler, L. Zhu and E.Gileadi, Electrochim. Acta, 48, 2551, (2003)] obtained high W deposit (˜67 at %) content Ni—W alloys from non-ammonia plating baths. However, current efficiency was reduced dramatically. The formation of a heteronuclear Ni—W monocitrato complex, [Ni(WO4)(cit)(H)]2− was proposed as to co-deposition of the Ni—W alloy.

d. IG-Mo Alloys

Holt and co-workers [L. E. Vaaler and M. L. Holt, Trans. Electrochem. Soc. 90, 43 (1946); L. E. Vaaler and M. L. Holt, Ibid., 94, 50 (1948); W. E. Clark and M. L. Holt, Ibid., 94, 244 (1948); M. H. Lietzke and M. L. Holt, Ibid., 94, 252 (1948); R. F. McElwee and M. L. Holt, J. Electrochem. Soc., 99 (2), 48 (1952).] extended their co-deposition studies of IG-W alloys to IG-Mo alloys from equivalent solutions. Hull Cells were initially used to determine operating conditions, such as pH, temperature, CD ranges for bright, metallic deposits. Similar to IG-W co-deposition, deposit Mo contents depended on the co-depositing IG metal with Fe>Co>Ni. Current efficiencies decreased with increased Mo content: Ni>Co>Fe. Complexation with citrate and tartrate were favored for Ni—Mo codeposition with malate (and malic acid) and glycolic acids for Co—Mo. Sodium citrate was superior to citric acid or ammonium citrate.

Landolt and Podlaha have published a series of papers [E. J. Podlaha and D. Landolt, J. Electrochem. Soc., 143, 885, (1996); E. J. Podlaha and D. Landolt, J. Electrochem. Soc., 143, 893, (1996); E. J. Podlaha and D. Landolt, Electrochem. Soc., 144, 1672, (1997)] on the codeposition of Ni-Mo alloys using rotating cylinder electrodes from the plating solutions of NH3, C6H5Na3O7.2H2O, Na2MoO4.2H2O, and NiSO4.6H2O. Alloy composition was affected by CD, electrode rotation rate, solution temperature, and species concentration, and Ni—Mo alloys of Mo content in excess of 50 wt % have been deposited [E. J. Podlaha and D. Landolt, J. Electrochem. Soc., 143, 885, (1996)]. For electrolytes of low Mo concentration in the present of excess Ni concentration, Mo content increased with increased electrode rotation rate and decreased with increased CDs. On the other hand, for electrolytes of low Ni concentration in the present of excess Mo concentration, Mo contents were independent of convection. A steady-state mathematical model was developed to predict the codeposition of Ni—Mo alloys [E. J. Podlaha and D. Landolt, J. Electrochem. Soc., 143, 893, (1996)]. This model assume that both Ni and Mo can complex with citrate in alkaline solutions, but the formation constant of Mo-citrate constant is much smaller than that of the Ni-citrate complex. The model predictions were in agreement with the observed trends in the experimental data [E. J. Podlaha and D. Landolt, J. Electrochem. Soc., 143, 885, (1996)]. Rotating cylinder electrodes of NiMo, CoMo, and FeMo alloys were electrodeposited for Mo solution concentrations much lower than the iron group species [E. J. Podlaha and D. Landolt, Electrochem. Soc., 144, 1672, (1997)]. Mo deposit content was higher in CoMo than NiMo and FeMo deposits due to a lower deposition rate of Co than Ni and Fe.

e. IG-V Alloys

Arcos et al. [C. Arcos, M. Schwartz and K. Nobe, AVG′, Electrochem. Soc., 95-15, 193, (1994); M. Schwartz, C. Arcos and K. Nobe, Plat. Surf Fin., 90, (6), 46, (2003)] reported the co-deposition (citrate baths) of binary and ternary alloys of the iron group metals and vanadium (i.e. Fe—V, Ni—V, Co—V, Co—Ni—V, Ni—Fe—V and Co—Fe—V) by DC and PC. Generally, vanadium deposit content increased with increase in pH (from 5.5 to 7.5) and increased CD (from 5 to 10 A/cm2). In binary alloys, at pH 7, V content decreased as: Fe>Ni>Co. Only Fe—V was obtained at pH below 7. For ternary alloys, the Co and V content increased and Fe decreased by increased pH for Co—Fe—V alloys. Convective mass transport and longer off time (PC) resulted in increased V deposit content. Co-deposited Co—Fe—V alloys had higher saturation magnetization compared to Ni—Fe (Perrnalloy type alloys) deposits for both DC and PC. The corrosion resistance of the deposits decreased as: Ni—Fe (DC)>Co—Fe—V (PC)>Co—Fe—V (DC)>Co—Fe (DC).

You et al. [B. Y. Yoo, M. Schwartz, and K. Nobe, Electrochim. Acta, 50, 4335, (2005)] investigated the electrodeposition of IG-V binary alloys from citrate solutions. Addition of NH3 and increasing pH lead to increase in V deposit content, but non-metallic deposits were obtained at solution pH>7. Increasing CD resulted in a linear decrease of V deposit content and a sharp decrease of current efficiency. In general, the V deposit content increased as follows: Ni (1 wt %)<Fe (2 wt %)<<Co (4 wt %).

5. Electrodeposition of Co-RE Alloys from Aqueous Solutions

Compared to non-aqueous investigations, few studies of the codeposition of Co—Sm alloys from aqueous solution have been reported. Most of these studied were done by the UCLA group [L. Chen, M. Schwartz, and K. Nobe, in Electrodeposited Thin Films, M. Paunovic and D. A. Scherson, Editors, PV 96-19, p. 239, The Electrochemical Society Proceedings Series, Pennington, N.J. (1996); M. Schwartz, F. He, N. Myung, and K. Nobe, in Magnetic Materials, Processes,and Devices V. Applications to Storage and Microelectromechanical Systems (MEMS), L. T. Romankiw, S. Krongelb, and C. H. Ahn, Editors, PV 98-20, p. 646, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999); N. V. Myung, M. Schwartz, and K. Nobe, in Fundamental Aspects of Electrochemical Deposition and Dissolution, M. Matlosz, D. Landolt, R. Aogaki, Y. Sato, and J. B. Talbot, Editors, PV 99-33, p. 263, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999); M. Schwartz, N. V. Myung, and K. Nobe, J. Electrochem. Soc., 151, C468, (2004).]. Zangari synthesized Sm—Co nanoparticles [J. Zhang, P. Evans, and G. Zangari, J. Magn. Magn. Mater., 283, 89, (2004).] by single short pulse electrodeposition from the solution proposed by Schwartz et al. [M. Schwartz, N. V. Myung, and K. Nobe, J. Electrochem. Soc., 151, C468, (2004).].

In 1996, Chen et al. [L. Chen, M. Schwartz, and K. Nobe, in Electrodeposited Thin Films, M. Paunovic and D. A. Scherson, Editors, PV 96-19, p. 239, The Electrochemical Society Proceedings Series, Pennington, N.J. (1996).] published the first study of IG-RE alloys from aqueous solution. Electrodepositions was carried out at room temperature and pH 4 from the plating solution containing RE mixtures, Co, Fe or Ni chloride salts, and various addition agents; soluble anodes were used. In direct current (DC) electrodeposition, it was noted that RE was not found in the deposits from the solutions of pH<4. RE deposit content (Co-RE and Ni-RE) increased with increasing current density (CD) from 5 to 20 mA/cm2. RE deposit content was higher for Fe-RE than Ni-RE and Co-RE. In pulsed current (PC) electrodeposition, higher temperatures and cobalt concentrations resulted in lower RE deposit content. For codeposition of metallic IG-RE alloys, specific addition agents (i.e. aminocarboxylates) in the plating solution were required.

In 1998, Schwartz et al. [M. Schwartz, F. He, N. Myung, and K. Nobe, in Magnetic Materials, Processes,and Devices V. Applications to Storage and Microelectromechanical Systems (MEMS), L. T. Romankiw, S. Krongelb, and C. H. Ahn, Editors, PV 98-20, p. 646, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999).] used plating solutions containing 0.3M RE metal ions (i.e. La, Ce, Nd, Gd and RE mixtures), 0.12M IG ions, 0.36M complexant (e.g., glycine, alanine and serine), 1M NH4Cl and 0.5M H3BO3 to obtain IG-RE alloys at room temperature. In DC electrodeposition, it was found that the addition of NH4Cl improved solution stability and deposit appearance. RE deposit content decreased in the order: glycine>serine>alanine. With glycine, the RE deposit content increased: Co<Fe<Ni. PC electrodeposition extended the effective peak CD range for metallic deposits. Crack density seemed to be directly related to the deposit RE content.

In 1999, Myung et al. [N. V. Myung, M. Schwartz, and K. Nobe, in Fundamental Aspects of Electrochemical Deposition and Dissolution, M. Matlosz, D. Landolt, R. Aogaki, Y. Sato, and J. B. Talbot, Editors, PV 99-33, p. 263, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999).] used chloride-based plating solutions containing 0.3M RE ions (i.e. Nd or Sm), 0.12M IG ions (i.e. Co and Ni), 0.36-0.72M glycine and 1M NH4Cl to obtain 10-RE alloys at room temperature. 0-0.4M DMAB (dimethylamineborane) was added to solutions to obtain ternary RE-IG-B alloys. Soluble IG served as anodes, and brass and stainless steel panels served as cathode substrates. Metallic Nd—Ni—B alloys were electrodeposited at pH6 and CD<40 mA/cm2. Increased CD led to increased Nd content, decreased B content and current efficiency. Increasing glycine/Ni ion concentration decreased Nd content and increased current efficiency. Metallic Sm—Co and Sm—Co—B alloys were obtained at pH 4-6.5, CD<40 mA/cm2. In the absence of DMAB in solutions, Sm content increased with increased CD; in the presence of DMAB, the opposite trend was observed. The crystal structures of Sm—Co and Sm—Co—B alloys were hexagonal closed pack (hcp) (CD<10 mA/cm2) or non-crystalline (CD>10 mA/cm2) Deposit grain size reduced from 128 to 38 nm by increased CD from 5 to 30 mA/cm2.

In 2004, Schwartz et al. [M. Schwartz, N. V. Myung, and K. Nobe, J. Electrochem. Soc., 151, C468, (2004).] used chloride- or sulfamate-based plating solutions containing 0.3 or 0.9M RE metal ions (i.e. Ce, Nd, Gd and Sm), 0.12M IG ions (i.e. Fe, Co and Ni), 0.36M complexant (e.g., glycine, alanine and serine), 1M NH4 ions (i.e. NH4Cl or NH4NH7SO3) to obtain IG-RE alloys at room temperature with soluble IG or insoluble Ti serving as anodes. The result agreed with previous studies [L. Chen, M. Schwartz, and K. Nobe, in Electrodeposited Thin Films, M. Paunovic and D. A. Scherson, Editors, PV 96-19, p. 239, The Electrochemical Society Proceedings Series, Pennington, N.J. (1996); M. Schwartz, F. He, N. Myung, and K. Nobe, in Magnetic Materials, Processes, and Devices V. Applications to Storage and Microelectromechanical Systems (MEMS), L. T. Romankiw, S. Krongelb, and C. H. Ahn, Editors, PV 98-20, p. 646, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999); N. V. Myung, M. Schwartz, and K. Nobe, in Fundamental Aspects of Electrochemical Deposition and Dissolution, M. Matlosz, D. Landolt, R. Aogaki, Y. Sato, and J. B. Talbot, Editors, PV 99-33, p. 263, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999).] that the RE deposit content increased with the increasing CD and solution pH. The Sm deposit content ranked as: Fe>Ni=Co. Metallic Co—Sm deposits did not extend beyond CD=400 mA/cm2 resulted in maximum Sm deposit content of 8 at %. Aminoacids were found to be effective complexing agents for the codeposition of RE alloys; glycine resulted in higher RE deposit contents than serine and alanine (glycine>serine>alanine) at room temperature. A mechanism for the codeposition of IG-RE alloys was proposed involving hetero-nuclear glycinato coordination complexes as a result of the zwitterionic characteristics of glycine. Surface adsorbed H atoms and/or direct electron transfer might result in step-wise reduction of the depositing metals.

In 2004, Zhang et al. [J. Zhang, P. Evans, and G. Zangari , J. Magn. Magn. Mater., 283, 89, (2004).] used the plating solution proposed by Schwartz et al. [N. V. Myung, M. Schwartz, and K. Nobe, in Fundamental Aspects of Electrochemical Deposition and Dissolution, M. Matlosz, D. Landolt, R. Aogaki, Y. Sato, and J. B. Talbot, Editors, PV 99-33, p. 263, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999).] to synthesize Sm—Co nanoparticles by single short pulse electrodeposition. Nanoparticle composition was a function of pulse amplitude (PCD 0.1-1.5 A/cm2) and pulse duration (Ton 5-100 ms); the relative atomic percent of Sm, defined as Sm/(Sm+Co), increased with increasing PCD and decreasing Ton XRD and XPS data indicate that hcp Co—Sm metallic alloys mixed with metal oxides have been obtained. The oxygen atomic ratio O/(Sm+Co+O) was a function of Ton. Increasing Ton decreased Sm content, while oxygen content increased up to a maximum of about 50 at %. For short Ton (few ms), oxygen content was as low as 3 at % (PCD=1000 mA/cm2). In-plane coercivities up to 5.3 kOe have been achieved for as-plated nanoparticles for Sm content of about 20 at %.

D. Applications of Magnetic Co—Sm Alloys

In one aspect, the disclosed method and compositions can be used in connection with high performance nanostructured permanent magnets including high temperature applications in aeronautical and aerospace applications. In a further aspect, the disclosed method and compositions can be used to produce dramatically miniaturized devices including electric motors, generators, actuators, alternators, gyros, magnetic couplings, magnetic bearings, centrifuges, hearing aid devices, computer hard drives, camcorders, industrial robots, maglev trains, and magnetic imaging systems (MIS).

In a yet further aspect, the disclosed method and compositions can be used to produce thick film (>1 nm) deposition for microelectromechanical systems (MEMS) devices.

In a still further aspect, the disclosed method and compositions can be used to produce ultra thin (<100 nm) controlled electrodeposition for nano-electromechanical systems, and nanosize biomedical devices and neuroelectrochemical applications.

E. Aqueous Electrodeposition Compositions

In one aspect, the invention relates to compositions for enhancing the aqueous electrodeposition of rare earth-transition metal alloys comprising: a water soluble salt of samarium, a water soluble salt of cobalt, and a comlexant. In a further aspect, the water soluble salt of samarium is samarium sulfamate. In a further aspect, the water soluble salt of cobalt is cobalt sulfate or cobalt sulfamate. In a further aspect, the composition comprises one or more supporting electrolytes. In a further aspect, the composition further comprises boric acid. In one aspect, the electrodeposition can be performed onto a conducting (e.g., metal) substrate.

In one aspect, the complexant can be one or more amino acid. The amino acid can be any amino acid known to those of skill in the art. In a further aspect, the amino acid is selected from amine carboxylates, for example, glycine, alanine, and serine,

In a further aspect, the complexant can be one or more hydroxycarboxylic acid. The hydroxycarboxylic acid can be any hydroxycarboxylic acid known to those of skill in the art. In a further aspect, the hydroxycarboxylic acid is selcted from malic, glycolic and lactic acids, citric, and tartaric acids.

In one aspect, the one or more supporting electrolytes (e.g., conducting salts) can be any electrolytes known to those of skill in the art. In a further aspect, the one or more electrolytes are selected from ammonium sulfamate, ammonium sulfate, ammonium chloride, and mixtures thereof.

In one aspect, the composition can comprise from about 0.25M to about 2.0M of the water soluble salt of samarium, from about 0.01M to about 0.5M of the water soluble salt of cobalt, from about 0.05M to about 0.5M of the complexant, and from about 0.1M to about 3M of the supporting electrolyte. In a further aspect, the composition can comprise 1M of the water soluble salt of samarium, 0.05M of the water soluble salt of cobalt, 0.15M of the complexant, and 1M of the supporting electrolyte.

In one aspect, the water soluble salt of samarium is samarium sulfamate. In one aspect, the water soluble salt of cobalt is cobalt sulfate or cobalt sulfamate. In one aspect, the complexant is an amino acid, for example, glycine. In one aspect, the complexant is a hydroxycarboxylic acid, for example, malic or citric acid. In one aspect, the conducting salt is ammonium sulfamate. In a further aspect, the water soluble salt of samarium is samarium sulfamate, the water soluble salt of cobalt is cobalt sulfate or cobalt sulfamate, the complexant is glycine, and the supporting electrolyte is ammonium sulfamate.

It is understood that the disclosed compositions can be used in connection with the disclosed methods.

F. Electrodeposition Methods

In one aspect, the invention relates to methods for electrodepositing a samarium-cobalt coating onto a conducting (e.g., metal) substrate, comprising placing an aqueous solution containing a water soluble salt of samarium, a water soluble salt of cobalt, one or more supporting electrolytes, and a comlexant into a plating bath, placing an anode and the substrate to be coated into the bath and connecting the anode and the substrate to a power supply, with the substrate acting as the cathode, adjusting the pH of the bath to a suitable operating level, and applying a direct current through the anode and substrate causing the samarium and the cobalt to migrate to, and adhere to, the substrate. In one aspect, the aqueous solution further comprises boric acid.

In one aspect, the method can further comprise an annealing step.

In one aspect, the complexant can be one or more amino acid. The amino acid can be any amino acid known to those of skill in the art. In a further aspect, the amino acid is selected from amine carboxylates, for example, glycine, alanine, and serine,

In a further aspect, the complexant can be one or more hydroxycarboxylic acid. The hydroxycarboxylic acid can be any hydroxycarboxylic acid known to those of skill in the art. In a further aspect, the hydroxycarboxylic acid is selcted from malic, glycolic and lactic acids, citric, and tartaric acids.

In one aspect, the one or more supporting electrolytes (e.g., conducting salts) can be any electrolytes known to those of skill in the art. In a further aspect, the one or more electrolytes are selected from ammonium sulfamate, ammonium sulfate, ammonium chloride, and mixtures thereof.

In one aspect, the aqueous solution can comprise from about 0.25M to about 2.0M of the water soluble salt of samarium, from about 0.01M to about 0.5M of the water soluble salt of cobalt, from about 0.05M to about 0.5M of the complexant, and from about 0.0001M to about 3M of the supporting electrolyte. In a further aspect, the aqueous solution can comprise about 1M of the water soluble salt of samarium, about 0.05M of the water soluble salt of cobalt, about 0.15M of the complexant, and about 1M of the supporting electrolyte.

In one aspect, the water soluble salt of samarium is samarium sulfamate. In one aspect, the water soluble salt of cobalt is cobalt sulfate or cobalt sulfamate. In one aspect, the complexant is an amino acid, for example, glycine. In one aspect, the complexant is a hydroxycarboxylic acid, for example, malic or citric acid. In one aspect, the supporting electrolyte is ammonium sulfamate. In a further aspect, the water soluble salt of samarium is samarium sulfamate, the water soluble salt of cobalt is cobalt sulfate or cobalt sulfamate, the complexant is glycine, and the supporting electrolyte is ammonium sulfamate.

In one aspect, a current density of from about 5 mA/cm2 to about 600 mA/cm2 is applied across the anode and cathode. In various aspects, the current density can be, for example, from about 5 mA/cm2 to about 300 mA/cm2, from about 5 mA/cm2 to about 100 mA/cm2, from about 5 mA/cm2 to about 50 mA/cm2, from about 5 mA/cm2 to about 20 mA/cm2, from about 10 mA/cm2 to about 300 mA/cm2, from about 10 mA/cm2 to about 100 mA/cm2, from about 10 mA/cm2 to about 50 mA/cm2, from about 10 mA/cm2 to about 20 mA/cm2, from about 0 mA/cm2 to about 300 mA/cm2, from about 0 mA/cm2 to about 100 mA/cm2, from about 0 mA/cm2 to about 50 mA/cm2, or from about 0 mA/cm2 to about 20 mA/cm2. In a further aspect, the DC current density is a DC current density. In a further aspect, the current is an alternating current. In a further aspect, the current is a pulsed current. In a further aspect, the current is applied with pulse current modifications varying with duty cycle and frequency.

In one aspect, the pH of the solution is from about 3 to about 6. In a further aspect, the pH of the solution is adjusted to from about 4 to about 6.5. In a further aspect, the pH of the solution is about 4.

In one aspect, the electrodeposition is conducted at a temperature of from grater than about 0° C. to less than about 100° C. In a further aspect, the electrodeposition is conducted at a temperature of from about 20° C. to about 80° C. In a further aspect, the electrodeposition is conducted at a temperature of from about 20° C. to about 60° C. In a further aspect, the electrodeposition is conducted at a temperature of from about 20° C. to about 40° C. In a further aspect, the solution temperature is from about 25° C. to about 60° C., for example, about 25° C., about 40° C., or about 60° C. In a yet further aspect, the electrodeposition is conducted at about room temperature.

The electrodeposition can be conducted with stirring. In a further aspect, the electrodeposition is conducted without stirring. In a further aspect, the electrodeposition is conducted with oscillatory stirring. In a further aspect, the electrodeposition is conducted with oscillatory stirring at a rate of from about 40 to about 60 cycles/min, for example, about 48 cycles/min.

It is understood that the disclosed compositions can be used in connection with the disclosed methods.

Also disclosed are samarium-cobalt coatings produced by the disclosed methods.

G. Nanostructured Magnetic Coatings

In one aspect, the invention relates to nanostructured magnetic coatings comprising a magnetic alloy of a rare earth metal and a transition metal. In a further aspect, the coatings are electrodeposited. In a further aspect, the coatings are electrodeposited from aqueous solution.

In one aspect, the rare earth metal is samarium. In one aspect, the transition metal is cobalt. In a further aspect, the rare earth metal is samarium and transition metal is cobalt. In a further aspect, the alloy comprises SmCo5 or Sm2Co17.

In certain aspects, the electrodeposited alloy contains sufficient samarium content to perform as a precursor to forming magnetic SmCo5 and/or Sm2Co17.

It is understood that the disclosed nanostructured magnetic coatings can be produced using the disclosed methods and compositions.

H. Experimental

The following examples are put forth so as to provide those of ordinary skill in the art with a complete disclosure and description of how the compounds, compositions, articles, devices and/or methods claimed herein are made and evaluated, and are intended to be purely exemplary of the invention and are not intended to limit the scope of what the inventors regard as their invention. Efforts have been made to ensure accuracy with respect to numbers (e.g., amounts, temperature, etc.), but some errors and deviations should be accounted for. Unless indicated otherwise, parts are parts by weight, temperature is in ° C. or is at ambient temperature, and pressure is at or near atmospheric.

1. Aqueous Electrodeposition of Rare Earth Metals and Transition Metals

Rare earth (e.g., samarium) and transition metal (e.g., cobalt) elements can be electroplated out of an aqueous solution to form bright metallic coatings on substrates by proper selection of the additives, such as complexing agent, solution pH, operating temperature, current density, complexing agent/metal ratio, complexing agent/transition metal ratio, and duty cycle. Particularly suitable complexing agents are glycine, alanine, and serine, which are all amino acids with a single carboxyl group. With the exception of cysteine, complexing agents evaluated which were not effective were amino acids with more than one carboxyl group or were not amino acids. Cysteine is an amino acid with one carboxyl group and a thio- group (—SH). The —SH apparently interfered with obtaining the desired result by causing the formation of hydroxides under the conditions evaluated.

While varying the operating conditions resulted in lesser concentrations of the desired materials in the films produced, conditions were still suitable for preparing RE containing coatings. The preferred complexing agent is glycine but other aminecarboxylates were also found to be effective. The preferred operating conditions include a current density of at least 5 mA/cm2, room temperature, a pH of 4, and a Co/glycine ratio of about 0.3. However, it has been found that addition of NH4Cl to the processing bath sharply reduced hydrogen evolution resulting in higher RE content deposits. Furthermore, while a pH of 4 is preferred, metallic deposits were obtained over a wide pH range including pH less than 4 and greater than 7. Stable alkaline plating baths for RE and TM salts are disclosed.

Plating solutions can be prepared containing various complexing agents, and transition metals (TM) (e.g., Co, Fe, Ni) and rare earth chloride salts. The solution pH can be adjusted upward with NaOH and lowered with HCl. Electrodeposition can be carried out at room temperature (RT) with DC current in the solutions containing TMCl2 and La, Ce, Nd and a rare earth mixture (MOLYCORP™) referred to below as the REM mixture. Other commercial rare earth mixtures are also suitable. The composition of the Molycorp™ mixture is given in Table 2.

TABLE 2
Rare Earth Mixture (Molycorp ™)
AnalysisEquivalentWt. Percent
Element% as oxide% as carbonateMetal
Ce1.01.30.7
La45.964.439.2
Nd12.918.011.1
Pr4.86.73.9
Sm0.40.60.3
Gd0.30.40.3
Y0.30.50.2
other RE~0.4~0.6~0.4
other elements~0.1~0.2

Primary test solutions were (A) Bath A—0.12M TMCl2, 0.5M B(OH)3, 0.36M complexing agent, 0.3 M RE or REM (B) Bath B—same as Bath A+1M NH4 Cl.

Solutions were either unstirred or stirred using a magnetic stirrer or by oscillatory stirring (48 cycles/min).

Each solution was used until accumulative exposure of 240-A-min/L at which point a new solution was prepared. The solution becomes less effective after 240-A-min/L because of consumption of the key ingredients in the rare earth mixture used. Brass or stainless steel panels were used as substrates. The substrates were mechanically cleaned and then subjected to a chemical treatment including soaking in alkaline cleaning solution for 10 min followed by rinsing with deionized water. Surfaces were then activated just before electrodeposition by immersion in 10% HCl for 30 sec. Soluble Co, Fe, or Ni anodes were used, depending on the solution, to minimize changes in the metal solution composition and to avoid known side effects due to insoluble anodes.

A Kraft Dynatronix power supply (model DRP 20-5) was used to provide pulse current (PC) waveforms and a PAR potentiostate/galvanostat (model 173) was used to provide DC current.

In order to evaluate the efficiency of the electrodeposition of RE-TM materials from solutions containing complexing agents, nitric acid was used to dissolve the deposited films. After evaporating the nitric acid solution to dryness, the resultant dried RE-TM residue was dissolved with deionized water and transferred to a plastic test tube. Hydrofluoric acid was added to separate the rare earths from ferrous metals by precipitation of rare earths fluorides. The precipitate was thoroughly washed with deionized water and transferred to a 50 milliliter beaker. Boric acid and nitric acid were then added to dissolve the precipitated rare earth fluorides. The solution was evaporated to dryness, resulting in water-soluble rare earth compounds. The dried sample was redissolved with deionized water and transferred into a 10 milliliter volumetric flask. One milliliter of ammonium acetate buffer and a complexing agent (alizarin red) were added. Ammonium acetate was used to buffer the solution to pH of 4.7 and the alizarin red was complexed with the rare earth to develop a specific color. After dilution to 10 milliliters, a spectrophotometer (λ=530 nm) was used to measure the absorbance. The absorbance obtained was then used to estimate the amount of rare earth in the deposit.

For plating solutions free from complexing agents, precipitation by oxalic acid was followed by dissolution of the oxalate precipitate with concentrated hydrocholoric acid, and finally precipitation with ammonia. The final white hydroxide precipitate from the ammoniacal solutions confirmed the presence of lanthanons in the deposit.

a. Effects of Complexing Agents

Using Bath A, eleven (11) complexing agents were investigated to study their effects on the production of RE-Co deposits and the stability of solutions. The solutions were stirred and exposed to current density of 20 mA/cm2 unless. The results are summarized in Table 3.

TABLE 3
Effects of Complexing Agent*
Rare Earth Content
Additivesin DepositAppearance
Glycine (REM)8.0%Bright metallic
Glycine (Ce)6.3%Grey metallic
Glycine (La)7.5%Black metallic
Glycine (Nd)3.4%Gray metallic
Alanine (REM)3.8%Bright metallic
Serine (REM)5.0%Bright metallic
Aspartic acid (REM)Not analyzedNon-metallic white
(RE) hydroxide
Glutamic acid (REM)Not analyzedNon-metallic white
(RE) hydroxide
Malic acid (REM)No REGrey metallic (pH 8.5)
Cysteine (REM)Not analyzedNon-metallic hydroxide
Glycolic acid (REM)0.2-1%Bright metallic
Lactic acid (REM)0.2-1%Bright metallic
EDTA (REM)No deposit
*Solutions (Bath A, pH 4) were stirred and electrodeposition was at 20 mA/cm2

It was found that the a-amino acids, glycine, alanine and serine stabilized the plating solution at pH 4, resulting in metallic deposits containing rare earths. The highest RE content in deposited films was obtained in solutions containing glycine while deposits of lower RE content were obtained with alaline and serine. All the deposits exhibited bright metallic appearance, which differed from the typical matte appearance of cobalt electrodeposits, indicating the effect of the rare earth elements. In order to test which element was preferentially deposited from the REM, separate runs were performed in the solutions containing glycine and Ce(Cl)3, Nd(Cl)3 or La(Cl)3. The presence of lanthanum in the solution gave a black metallic deposit containing 7.5% lanthanum, 3.4% Nd was obtained with NdCl3 and these Ce(Cl3) produced a gray metallic deposit with a 6.4% Ce in the films. In these cases, the 3 RE content of the deposit was lower than that when the RE mixture (8%) was used.

Solutions containing aspartic acid and glutamic acid were not stable and produced uniform white precipitates which consisted of RE hydroxides instead of metal films. Black deposits were obtained from the solutions containing cysteine and those cysteine solutions were also not stable.

The solutions containing glycolic acid or lactic acid were cloudy at pH4 due to the formation of small amounts of hydroxides. However, bright metallic deposits containing small amounts of RE were obtained from filtered solutions. EDTA formed strong complexes with Co. As a result, no deposits were obtained from the EDTA containing solutions.

In addition to the results shown in Table 3, a Nd—Ni deposit of 6% Nd was obtained from Bath B using ethylene diamine as a complexing agent. The solution (pH5) was unstirred and deposits were obtained at 15 mA/cm2.

b. Effect of Direct Current Deposition

To evaluate the effect of current density on resultant deposits, electrodeposition was carried out at room temperature and current densities of 5, 10, and 20 mA/cm2 for Co-RE, Ni-RE and Fe-RE solutions containing glycine at pH4. The solution contained 0.12M (Fe, Ni, Co) Cl2, 0.5M B(OH)3, 0.36M glycine and 0.3 RE (La, Ce, Nd), or REM. FIG. 11 compares the dependence of the rare earth content (% rare earth) of the deposited films at different current densities. Generally, the percentage of rare earth in the film increased with increasing current density. Deposit content of the rare earths were greater in Ni alloys, less in Fe alloys and least in Co alloys. Rare earth deposit contents are typically greatest from unstirred solutions, a lesser amount from solutions mixed by oscillatory stirring and least from more vigorous agitation with a magnetic stirrer. Thus, mass transfer effects can be important in the efficacy of RE-TM electrodeposition.

C. Effect of Temperature

Electrodeposition from magnetic stirred Bath A containing CoCl2 and the rare earth mixture (REM) was run at both room temperature and 65° C. to examine the temperature dependence of Re—Co deposits. It was found that at the same current density (20 mA/cm2), the rare earth in the deposits at 65° C. was ˜3%, which was less than half the 6.6% obtained at room temperature. Thus, the cobalt deposition rate is greatly enhanced and the RE deposition reduced as temperature is increased. In other words, a lower temperature during electrodeposition favors RE deposition.

d. Effect of Complexing Agent to Metal Ratio

The ratio of the glycine concentration to metal concentrations in magnetic stirred solutions also had a measurable effect on RE-Co electrodeposition. FIG. 12 shows the effects of glycine/Co solution ratios with CoCl2 held constant at 0.12M on the deposit RE content obtained at room temperature with a current density of 20 mA/cm2 and a pH of 4. There appears to be a plateau or an approach to a maximum in deposited RE content as the glycine/Co ratio approached 1. At glycine/Co ratio>1, a sharp decrease in the deposit RE content with increasing ratios was observed (FIG. 12).

e. Effect of Co(Cl)2+Glycine

In this study, the magnetic stirred solution RE concentration was maintained constant at 0.3M while the combined concentrations of Co(Cl)2+glycine was increased at a constant ratio: 1Co:3glycine. FIG. 13 shows increased Co(Cl)2+glycine concentrations resulted in decreased deposit RE content. At a combined total concentrations of 1.5M, practically no RE was deposited indicating the possible inhibitory-effect of increasing addition agent concentrations. Again, operating conditions were room temperature, pH of 4 and a current density of 20 mA/cm2.

The duty cycle for PC electrodeposition is defined as ton/(ton+toff), and the average current density is the peak current density times the duty cycle. Pulsed current deposition of RE-Co alloys was performed at an average current density of 20 mA/cm2 with Ton at 5 msec. FIG. 14 shows that the deposit RE content was fairly constant at ˜4.5±5% at duty cycles from 0.1 to 0.8. In this range, the peak cathodic current densities ranged from 200 to 25 mA/cm2, along with decreasing off-times of 45 to 1.75 msec, respectively. At duty cycles greater than 0.8, approaching DC plating, the deposit RE content increased to ˜6.+1% and was similar to that obtained with constant DC current.

As the peak cathodic current density increased, the required longer off-times (relaxation times) permitted sufficient diffusion of either or both the Co or RE species into the cathode diffusion layer. However, at any peak cathodic current density greater than DC, the diffusion of the RE was insufficient to provide the necessary replenishment, resulting in lower deposit content, although the bulk solution concentration was three times that of cobalt. More Co deposited during the on-time indicating either fast deposition rates or mass transfer compared to the RE.

For Co—Re deposition, deposit RE content was relatively constant with PC deposition up to duty cycle of 0.8 and then increased at higher duty cycle. DC electrodeposition gave the highest amount of RE in the films. Temperatures greater than room temperature increased additive to metal ratio, and increased cobalt concentration resulted in lower RE in the films.

f. Effect of Solution Ph and NH4Cl

The solution pH can be important to the electrodeposition process. The pH can affect the onset of the hydrogen evolution reaction, the composition of the deposits, the current efficiencies and the stability of the solution. Addition of NH4Cl to Bath A was an effort to lessen the rate of hydrogen evolution. FIG. 15 illustrates the interdependence of current density with solution pH on the composition of deposits obtained from TM-Nd-glycine solutions. In general, the deposit Nd content increased fairly linearly with increasing current density and increasing solution pH in the range of 5-40 mA/cm2 and pH4-5.4, respectively, the exception being Nd—Ni deposits which exhibited a maximum deposit content at 10 mA/cm2 and solution pH of 4.8.

It was observed that the presence of NH4Cl significantly decreased hydrogen evolution during electrodeposition of RE-TM alloys. As a result the pH range to obtain metallic deposits was increased. For example, 29% Ce in Ce—Ni deposits were obtained with glycine @ pH2.7 and 15 mA/sq.cm (Bath B) and 23% Nd was obtained in Nd—Ni deposits with alanine)@pH7 and 20 mA/sq.cm (Bath B). Furthermore, deposit RE content was generally higher in solutions containing NH4Cl. For example, for Ce—Ni deposits at 5 and 20 mA/cm2 with oscillatory stirring (Bath B), Ce contents were 10.5% and 22.5%, respectively. In comparison 8.2% and 16.2% were obtained from Bath A.

g. Mass Transfer Effects

The degree of solution agitation during electrodeposition of RE-TM alloys can have an effect on the RE content of the deposits. FIG. 16 shows that the Ce content in Ce—Ni deposits was less from oscillatory stirred solutions (48 cycles/min) compared to unstirred solutions. Further, RE deposit contents were even lower from solutions agitated more vigorously using a magnetic stirrer. On the other hand, visual inspection of the deposits indicates that solution agitation improved the quality (appearance) of the deposits. For the electrodeposition of bright metallic or ferrous metal-RE alloys, the most effective complexing agents appear to include glycine, alanine and serine. These complexing agents are amino acids with a specific chemical structure, namely a single carboxyl group and thus differ chemically from the other sampled complexing agents which were not found to be suitable. Therefore, it would appear that other amino acids with single carboxyl groups would be suitable compounds to create the same result under similar operating conditions and solution compositions. Other types of complexing agents investigated were either not as effective or ineffective, usually resulted in precipitation of hydroxide in the solution and/or in the deposited films or prevented deposition of the RE or resulted in unacceptable appearing films.

2. Hull Cell Studies

Co—Sm permanent magnets, such as Sm2Co17 and SmCo5, require 10.53 and 16.67 at % Sm content, respectively. To satisfy the composition requirements of Co—Sm magnets, the alloys produced by electrodeposition must contain enough Sm content. Therefore, high Sm content Co—Sm alloys electrodeposited from aqueous solution is the initial goal of this research.

An electrodeposition process can be operated successfully only when the key parameters are properly controlled. These are components and compositions of plating baths (e.g. metal ions, supporting electrolytes and additives) and operating conditions (e.g. current density (CD), solution temperature, pH, fluid dynamics and current waveforms). To obtain high Sm deposit content, these parameters need to be studied carefully.

The Hull cell is an effective screening device often used by electroplaters to solve problems of the electroplating process. The Hull cell has been recognized as a powerful tool to study the approximate deposit properties. Generally, the Hull cell provides information regarding the deposit characteristics over a wide range of CDs and multiple experimental results in a single experiment. For its high efficiency, Hull cell technology was chosen to determine the dependence of Sm deposit content on deposit parameters and coupling between deposit parameters in the electrodeposition of Co—Sm alloys. Although the result by the Hull cell is less accurate and more limited than by parallel electrodes, it still provides a good approximation to the trends in Sm deposit content by varying the electrodeposition parameters in the initial investigations of the electrodeposition of Co—Sm alloys.

FIG. 17 shows the flowchart of a Hull cell experiment which mainly includes four parts: pretreatment of cathode, DC or PC electrodeposition, post-treatment of specimen and characterization. Fundamentals, definitions, experimental setup, design of Hull cell study, pretreatment and post-treatment, and characterization and analysis of the specimens will be described in the following discussion.

a. Fundamentals of the Hull Cell

The Hull cell, developed by R. O. Hull [R. O. Hull, U.S. Pat. No. 2,149,344 (1939)], is a miniature trapezoidal plating cell (267 mL vol) which provides a current density (CD) range on the cathode test panel, depending on the applied current. FIG. 18 shows Hull cell cathode test panel at the right end (point b), having the longest cathode-anode distance (Db), resulting in the lowest CD; the CD continuously increases as the current path along the cathode decreases and reaches a maximum at point a, the shortest cathode-anode distance (Da).

Thus, it is universally used as an economical screening device to evaluate the effects of solution compositions and applied operating conditions on the deposit appearance, composition and crystal structure as a result of incremental CDs on a single cathode surface, especially for initial investigations of alloy electrodeposition. Further, deposits on selected portions of the test panel can be analyzed by energy dispersive spectroscopy (EDS) and X-ray diffraction (XRD) to provide additional information regarding compositions and crystal structures. The current density at which the deposit no longer has a metallic appearance, referred to as “burnt” by practicing electroplaters, was defined as the maximum current density (CDmax) for a particular plating system.

To minimize solution concentration and temperature gradients during electrodeposition, the Hull cell was equipped with a motorized slide mechanism with an attached paddle located alongside the cathode providing a reciprocal horizontal motion (agitation) with a sweep rate of 80 cycles/min, regulated by a variable resistor.

Hull developed an equation describing the CD distribution on the test panel for a typical 267 mL Hull cell:


CD=I(27.7−48.7 log L)A/ft2 (ASF) (Equation 6)

where I is the applied current in amperes, and L is the position on the test panel in inches from the low CD end (point b). This equation was derived basing on the standard 10×5 cm (area=50 cm2) Hull cell panel. The design of the Hull cell allows us to obtain deposits at different CDs on a single substrate. In other words, it saves a lot of time by providing a spectrum (Hull cell pattern) of a deposit obtained at continuous changing CD along its length on a single panel. The deposits can then be analyzed to determine properties at a particular CD.

However, the design of the Hull cell is based on primary current density distribution neglecting the secondary density distribution [L. J. Durney, Electroplating Engineering Handbook (4th edition), Van Nostrand Reinhold Company Inc., Taiwan, (1984), pp. 461-473] due to the depletion of metal ions at cathode surface during the electrodeposition. The apparent CDs on the test panel, as calculated with Equation 6, only provide an approximation of the “true” CDs. Therefore, the analysis in Hull cell only shows the trends rather than the precise values of deposit properties changed by electrodeposition parameters. Only Sm deposit contents (by EDS) and crystal structures (by XRD) were analyzed in the Hull cell test. In addition, analysis of the deposit at CDmax is difficult to determine because cutting a piece of specimen at CDmax non-metallic deposits at CDs slightly higher than CDmax were included. Compared to the result by parallel electrodes depositions, the Hull cell result is less accurate and more limited. Although the Hull cell has these drawbacks, it is still an efficient device providing a good approximation of trends for the preliminary investigation of the electrodeposition of Co—Sm alloys.

b. Definitions and Parameters

DC&PC electrodeposition: Metallic deposit is defined as a deposit with a visual metallic-appearance, Sm content (at %) is defined as the atomic percentage of Sm in deposited total metals, Sm content

(at%)=Smtotalmetals(at%)=SmSm+Co(at%)

DC electrodeposition: CD is the current density defined as current per unit deposit area, CDmax, is defined as the highest CD to obtain metallic-appearing deposits.

PC electrodeposition: PCD is peak current density defined as the maximum CD in one complete pulse cycle, PCDmax is defined as the highest PCD to obtain metallic appearing deposits, Ton, is the time duration of the on-current in one complete pulse cycle, Toff is the time duration of the off-current in one complete pulse cycle, Period is the time duration for one complete cycle, period=Ttotal=Ton+Toff, Frequency (f) is defined as the number of complete cycles per second,

f=1period=1Ton+Toff

Duty cycle (γ) is defined as the ratio of Ton to period,

γ=TonTon+Toff=Tonf

c. Experimental Setup and Design

The setup for Hull cell electrodeposition is shown in FIG. 20. A Kraft Dynatronix power generator (model DRP 20-5-10) served as power source to supply current needed for DC and PC electrodeposition, a coulometer to measure the total charge passed, and an oscilloscope to monitor the waveform during pulse current electrodeposition. Masked brass panels (10×5 cm) with exposed 15 cm2 (10×1.5 cm) deposit area for DC and with 7.5 cm2 (10×0.75 cm) for PC electrodeposition served as cathodes and a platinum sheet (5×5 cm) was used as the anode.

The reduction of deposit area from 50 cm2 (whole brass panel) to 15 cm2 (DC) or 7.5 cm2 (PC) increased the CD range for the parametric studies. Therefore, equation 1 was modified by using 15 cm2 (or 7.5 cm2) deposit area:


CD=I(85.8.−150.8 log L), mA/cm2 (4.5 A/15 cm2) (Equation 2)


CD=1(171.6.−301.610 log L), mA/cm2(4.5 A/7.5 cm2) (Equation 3)

where 1 is the total applied current in amperes, and L is the position on the test panel in inches from the low CD end (right end).

The deposits were obtained in a 267 mL Hull cell filled with the plating baths as shown in Table 4.

TABLE 4
Plating baths used in Hull cell studies
Bath #Sm sulfamateCo sulfateGlycineNH4 SulfamatepH
11 M0.05 M0.15 M5.7
21 M5.8
31 M0.15 M5.7
40.05 M5.6
50.05 M0.15 M4.5
61 M0.05 M5.9
71 M0.05 M  3 M4.0
81 M0.05 M0.15 M1 M5.9
*The pH values of the plating baths were measured at 25° C.

Bath 1 was used to study the effect of CD and temperature; baths 2 and 3 were used to study the effect of glycine on the formation of Sm oxide and hydroxide; baths 4 and 5 were used to study the effects of glycine on the electrodeposition of Co; baths 1, 6 and 7 were used to study the effect of the glycine concentration on the codeposition of Sm and Co; baths 1 and 8 were used to study the effect of ammonium sulfamate, as the supporting electrolyte.

Unless otherwise noted, the total charge passed was 100 coulombs for DC and 50 coulombs for PC electrodeposition to provide deposits thick enough to be analyzed. The deposit area was 15 cm2 for DC and 7.5 cm2 for PC electrodeposition. The applied charge density remained constant as 6.67 C/cm2 for both DC and PC. The applied current was 4.5 A for DC (25 and 60° C.) and for PC at 25° C. and 7A at 60° C. providing a wide range of CDs. Solutions were not agitated during electrodeposition.

(1) Pretreatment and Post-Treatment

Before electrodeposition, the brass panels were mechanically cleaned with a brush, soaked in 0.1M NaOH for 10 min., rinsed in deionized water, immersed in 10% HCl for 30 seconds and then rinsed with deionized water.

After the Co—Sm alloy was deposited for 100 coulombs in DC or 50 coulombs in PC, deposits were removed from the plating solution, rinsed with deionized water, and then dried with nitrogen gas. Disk-shaped specimens of diameter of 3.2 mm (specimen area=8.04 mm2) were die-punched out from deposits for analysis.

(2) Characterization and Analysis

The main purpose of the Hull cell study is to determine the trend in Sm deposit content by varying the electrodeposition parameters; the Sm deposit content was determined by energy dispersive x-ray spectroscopy (EDS) by a Kevex detector within a Cambridge scanning electron microscopy (SEM) (model Stereoscan 250). A PANalytical X-ray diffraction (XRD) (model X'Pert Pro) was used to examine crystal structures of deposits by {tilde over (Θ)}2Θ scan method. Unless otherwise noted, the experimental data presented are restricted to metallic-appearance deposits.

(3) Energy Dispersive X-Ray Spectrometer (EDS)

EDS measures the energy and intensity distribution of X-rays generated by the bombardment of electron beam on the specimen. The composition of the specimen can be obtained by comparing the peak intensities of Co (Kα1, 6.93 ev) and Sm (Lα1, 5.62 ev) to the intensities of the internal standard of pure Co and Sm, respectively, to get the k′-ratios

(k=IspecimeniIpureelementi)

then calibrated by ZKF method (a matrix correction technology) to get the k-ratios of Co and Sm. K-ratios, which are proportional to the weight percent of elements in the specimens, were used to obtain the weight and atomic percent of Co and Sm. An example of energy dispersive spectrum of an electrodeposited Co—Sm alloy is given in FIG. 21 and the analysis result is shown in Table 5. The elemental composition in a defined scan area can be easily determined to a high degree of precision (−0.1 wt. %).

TABLE 5
Compositional analysis results of an electrodeposited Co—Sm alloy
ElementPeakPeak Intensity (cps)K-RatioWeight %Atomic %
Co11793.40.855885.7793.89
Sm1149.20.144214.236.11

(4) X-Ray Diffraction (XRD)

XRD [L. V. Azaroff, Elements of X-ray crystallography, McGraw-Hill, N.Y., (1968)] is a technique in crystallography which can be used to determine the crystal structures of the specimen by characterizing its diffraction pattern with Bragg's law. The shape and size of the unit cell determines the angular position (2Θ) of the diffraction lines; the arrangement of the atoms within the unit cell determines the relative intensities of the lines. Information regarding states of Co—Sm alloys (e.g. crystalline, non-crystalline, intermetallic compound), non-metallic compounds (e.g., oxide and hydroxide), and prefer orientation (PO) of deposits can be examined by diffraction peaks. These characteristics of electrodeposited Co—Sm alloys controlled by electrodeposition parameters are very useful to study magnetic, electric and mechanical properties of deposits.

The grain size of the crystallites in out-of-plane direction (perpendicular to film plane) can be estimated from the measured width of their diffraction peaks by Scherrer's formula [B. D. Cullity and S. R. Stock, Elements of X-ray diffraction (3rd edition), Prentice Hall, N.J., (2001), p170]:

t=0.9λBCos(θB)

where λ is the wavelength of X-ray used to obtain the diffraction pattern; B is the full-width at half maximum (FWHM), and ΘB is the Bragg's angle of the diffraction peak.

(5) Effect of Applied Charge, Current and Deposit Area

As indicated, various applied charges (100 C/15 cm2 for DC and 50 C/7.5 cm2 for PC) and currents (4.5 A/15 cm2 for DC (25 and 60° C.), 4.5 A/7.5 cm2 for PC at 25° C. and 7 A/7.5 cm2 for PC at 60° C.) were applied to obtain Hull cell patterns. To confirm the results of CDmax (or PCDmax) obtained at different operating conditions can be compared, some pre-tests were done as follows.

Effect of different applied currents: The purpose of this test was to evaluate the consistency of PCDmax at various applied currents. FIG. 22 shows the Hull cell patterns obtained at 60° C. from bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine) by PC electrodeposition of applied current of 4.5 and 7 A. The PCDmax, of 4.5 A was about 990 mA/cm2 which was close to the PCDmax of 7 A (about 1000 mA/cm2). It was concluded that PCDmax at different applied currents varied little and the Hull cell patterns obtained at different applied currents (4.5 and 7 A) can be compared. On the other hand, the application of current 7 A provided a wider PCD range on Hull cell patterns obtaining more information of the oxide/hydroxide region (FIG. 22(b) compared to 4.5 A (FIG. 22(a)). Therefore, a current of 7 A was applied in PC at 60° C. instead of 4.5 A for a wider PCD range. PCDmax remained unchanged, and Hull cell patterns obtained at the current of 4.5 and 7 A can be compared.

Effect of different applied charges: The purpose of this test was to check the consistency of PCDmax at a fixed charge density. FIG. 23 shows the Hull cell patterns by PC electrodeposition at 25° C. from bath 1 for applied charge/deposit area of 100 C/15 cm2 and 50 C/7.5 cm2 (charge density was fixed at 6.67 C/cm2). The PCDmax of 100 C/15 cm2 was about 190 mA/cm2 which was close to the PCDmax of 50 C/7.5 cm2 (about 200 mA/cm2). It was concluded that as long as the charge density was constant (charge/deposit area=6.67 C/cm2), PCDmax varied little. Therefore, in PC electrodeposition, the PCDmax of Hull cell patterns obtained at 100 C/15 cm2 and 50/7.5 cm2 can be compared.

(6) Results and Discussions of DC Electrodeposition

Effect of current density and solution temperature: Bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine) was used to study the effects of current density and solution temperature on the electrodeposited Co—Sm alloys. The experimental conditions are shown in Table 6, and their Hull cell patterns in FIG. 24.

TABLE 6
The effects of current density and solution temperature
[Sm(NH2SO3)3][CoSO4]TCDmax
EXP #Bath(M)(M)[Glycine] (M)pH(° C.)(mA/cm2)
311.000.050.155.72550
660750
4580850
* Total charge = 100 C, applied current = 4.5 A, substrate area = 15 cm2, pH = 5.7, no agitation.

Deposits obtained from bath 1 at 25° C., for example, showed metallic-appearing for CD below 50 mA/cm2, burnt between 50 and 100 mA/cm2, and white powder at CD above 100 mA/cm2 (FIG. 24(a)). To characterize these regions, XRD was used to study their phase compositions. From the result of XRD patterns of deposit #3 (FIG. 25), the metallic region was non-crystalline and contained a weak diffraction peak of Sm(OH)3 (FIG. 25(a)). The burnt region exhibited not only Sm(OH)3 but also Co(OH)2 and SmO peaks (FIG. 25(b)). The non-metallic white powder region (FIG. 25(c)) contained Sm(OH)3, Co(OH)2 and mixtures of Sm and Co oxides.

Generally, metallic deposits could be obtained only below critical CD; at higher CDs, non-metallic appearing deposits containing hydroxides and oxides were obtained. This critical CD was defined as the maximum current density (CDmax) to obtain metallic deposits. It was observed that CDmax increased with increasing solution temperature (Table 6 and FIG. 24). For example, CDmax increased from 50 to 850 mA/cm2 by increased solution temperature from 25 to 80° C. Higher CD resulted in higher Sm deposit content as shown in FIG. 26. Therefore, high Sm deposit content of 25 at % can be obtained from bath 1 at 60° C. and 650 mA/cm2 exceeding the composition requirement of 16.67 at % for SmCo5. (Note: The deposit obtained at 60° C. and 750 mA/cm2 (CDmax) should have higher Sm deposit content than at 650 A/cm2. However, it was difficult to measure the Sm content at CDmax because it was too close to the non-metallic region.)

(7) Effect of Fluid Dynamics

Bath 1 was used to study the effect of fluid dynamics on electrodeposition of Co—Sm alloy. Solution agitation was achieved by a periodic reciprocal movement of the paddle along the Hull cell panel controlled by a motor (FIG. 27).

The experimental conditions are provided in Table 7, and the Hull cell patterns of the deposits are shown in FIG. 28. Solution agitation didn't significantly affect CDmax (FIG. 28) and Sm deposit content (FIG. 29) at either 25 or 60° C. To study mass transfer effect in the electrodeposition of Co—Sm alloy, a more controlled method, RDE (rotating disk electrode), was used, and the results are discussed in DC electrodeposition studies.

TABLE 7
The effect of fluid dynamics
[Sm(NH2SO3)3][CoSO4][Glycine]AgitationTCDmax
EXP #Bath(M)(M)(M)(cycles/min)(° C.)(mA/cm2)
311.000.050.1502550
98050
6060750
1280750
*Total apply charge = 100 C, applied current = 4.5 A, substrate area = 15 cm2 (10 cm × 1.5 cm), pH = 5.7.

(8) Effect of Glycine on the Electrodeposition of Sm and Co

Like other RE metals, metallic Sm has not been deposited from aqueous solution, generally attributed to its very negative reduction potential [W. M. Latimer, The Oxidation States of the Elements and Their Potentials in Aqueous Solution, Prentice-Hall, N.Y., pp. 286-295 (1952)] (E°<−2.3VSHE). Compared to Sm, water has a much less negative reduction potential [W. M. Latimer, The Oxidation States of the Elements and Their Potentials in Aqueous Solution, Prentice-Hall, N.Y., pp. 29-37 (1952)] (2H2O+4e→2OH+H2, E≧−0.826V). Typically, hydroxyl ions generated by hydrogen evolution react with Sm ions to form hydroxides. However, metallic codeposits of Co—Sm have been obtained from aqueous solutions containing glycine and its derivatives [L. Chen, M. Schwartz, and K. Nobe, in Electrodeposited Thin Films, M. Paunovic; M. Schwartz, F. He, N. Myung, and K. Nobe, in Magnetic Materials, Processes, and Devices V. Applications to Storage and Microelectromechanical Systems (MEMS), L. T. Romankiw, S. Krongelb, and C. H. Ahn, Editors, PV 98029, p. 646, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999); N. V. Myung, M. Schwartz, and K. Nobe, in Fundamental Aspects of Electrochemical Deposition and Dissolution, M. Matlosz, D. Landolt, R. Aogaki, Y. Sato, and J . B. Talbot, Editors, PV 99-33, p. 263, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999).; M. Schwartz, N. V. Myung, and K. Nobe, J. Electrochem. Soc., 151, C468, (2004).]. Therefore, it is of interest to study the effect of glycine on the codeposition of Co—Sm alloys.

Various solutions were used to study the effect of glycine on the formation of samarium oxide and hydroxide (bath 2 and 3), the electrodeposition of cobalt (bath 4 and 5), and the codeposition of samarium and cobalt (bath 1, 6 and 7).

(9) Formation of Sm Oxide and Hydroxide

Bath 2 (1M Sm sulfamate) and bath 3 (1M Sm sulfamate, 0.15M glycine) shown in Table 8 were used to study the effect of glycine on the formation of Sm oxide/hydroxide at 25 and 60° C. FIG. 30 shows the results of the Hull cell patterns.

TABLE 8
The effect of glycine on formation of Sm oxide/hydroxide
EXP[CoSO4][Glycine]T
#Bath[Sm(NH2SO3)3] (M)(M)(M)pH(° C.)
4121.00005.825
4260
4330.155.725
4460
*Total charge = 100 C, applied current = 4.5 A, substrate area = 15 cm2, no agitation.

Addition of glycine to Sm sulfamate solution did not result in metallic Sm deposits. However, it seems to play a role to stabilize Sm ions in solution and reduce the formation of hydroxides and oxides in deposits. It was noted that a mixture of Sm hydroxide/oxide was formed when CD reached a critical point; addition of glycine to the electrolyte decreased the formation of hydroxides and oxides by extending this critical CD. At 25° C., for example, the critical CD increased from 60 to 150 mA/cm2 by addition of 0.15M glycine to 1M Sm sulfamate (FIGS. 30(a) & (b)). Similar results were found at 60° C. (FIGS. 30(c) & (d)). For solutions with glycine present, higher CDs were needed to form Sm hydroxide. It has been reported that glycine derivatives complex Sm ions [J. Torres, C. Kremer, E. Kremer, H. Pardo, L. Suescun, A. Mombru, S. Dominguez and A. Mederos , Inorg. Chinn. Acta, 355, 442 (2003); J. Torres, C. Kremer, E. Kremer, H. Pardo, L. Suescun, A. Mombru, S. Dominguez and A. Mederos , J. Alloy Comp., 323-324, 119 (2001)] and prevent the precipitation of Sm(OH)3 [F. Medrano, A. Calderon and A. K. Yatsimirsky, Chem. Commmun., 1968, (2003)] by complexation of glycine and Sm ions reducing reaction of Sm3+ and hydroxyl ions.

It was also observed that increasing solution temperature also depressed formation of Sin hydroxide and oxide without (FIGS. 30(a) & (c)) or with glycine (FIGS. 30(b) & (d)) by extending the critical CD. In brief, it was found that addition of glycine and an increase in solution temperature depressed the formation of Sm oxide and hydroxide in the deposits.

(10) Electrodeposition of CO

Bath 4 (0.05M Co sulfate) and bath 5 (0.05M Co sulfate, 0.15M glycine) shown in Table 9 were used to study the effect of glycine on the electrodeposition of Co at 25 and 60° C. FIG. 31 shows the Hull cell patterns. Adding glycine (bath 5) effectively increased CDmax, especially at 60° C. CDmax increased from 90 to 500 mA/cm2 at 60° C. by the addition of 0.15M glycine into bath 4. It has been reported that the Co-glycine complex can inhibit the formation of Co(OH)2.[C. F. Diven,F. Wang, A. M. Abukhdeir, W. Salah, B. T. Layden, C. F. Geraldes, and D. M. Freitas, Inorg. Chem., 42, 2774, (2003)] Therefore, addition of glycine appears to prevent the formation of Co(OH)2, and extended the metallic deposit region to higher CDs.

TABLE 9
The effect of glycine on electrodeposition of Co
[Sm(NH2SO3)3][Glycine]CDmax
EXP #Bath(M)[CoSO4] (M)(M)pHT (° C.)(mA/cm2)
151400.0505.62530
1526090
15350.154.525110
15460500
*Total charge = 100 C, applied current = 4.5 A, substrate area = 15 cm2 (10 cm × 1.5 cm), no agitation.

An increase in solution temperature also increased CDmax. Deposits obtained from bath 5 increased CDma, from 110 to 500 mA/cm2 by increasing solution temperature from 25 to 60° C. In summary, addition of glycine and increase of solution temperatures resulted in higher CDmax.

(11) Electrodeposition of Co—Sm Alloys

As discussed in the previous sections, glycine can form complexes individually with both Sm and Co ions. Sm3+ complexed with glycine can not be electrodeposited to Sm (FIG. 30). Previous work [M. Schwartz, F. He, N. Myung, and K. Nobe, in Magnetic Materials, Processes, and Devices V. Applications to Storage and Microelectromechanical Systems (MEMS), L. T. Romankiw, S. Krongelb, and C. H. Ahn, Editors, PV 98029, p. 646, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999)] has shown that Co—Sm alloys can be electrodeposited from aqueous solutions containing Sm3+, Co and glycine (or other appropriate complexers). Electrodeposition of Co—Sm alloys have been studied at 25 and 60° C. in the absence and the presence of glycine at two concentrations; Bath 6 (1M Sm sulfamate, 0.05M Co sulfate), bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine) and bath 7 (1M Sm sulfamate, 0.05M Co sulfate, 3M glycine) were selected, as shown in Table 10. The Hull cell patterns and Sm content of deposits from these three baths are shown in FIG. 32 (for 25 and 60° C.) and FIG. 33 (for 60° C. only), respectively. In the absence of glycine, deposits had metallic appearance but contained considerable hydroxides/oxides.

TABLE 10
The effect of glycine on electrodeposition of Co—Sm alloys at 25 and 60° C.
[CoSO4][Glycine]TCDmax
EXP #Bath[Sm(NH2SO3)3] (M)(M)(M)pH(° C.)(mA/cm2)
1561.000.0505.82520
310.155.750
2573.004.040
2261.000.0505.860150
610.155.7750
3173.004.0650
*Total charge = 100 C, applied current = 4.5 A, substrate area = 15 cm2, no agitation.

The addition of glycine extended the metallic deposit region by increasing CDmax. CDmax from 0 to 3M (FIG. 32) increased, reached a maximum, and then decreased with increased glycine concentration. At 60° C., for example, the CDmax increased from 150 mA/cm2 (no glycine), reached a maximum of 750 mA/cm2 (0.15M glycine), then slightly decreased to 650 mA/cm2 (3M glycine).

Differing from other metallic deposits observed in previous sections, it was noted that the surfaces of “metallic” deposits from bath 6 (no glycine) at 25 and 60° C. were covered by a thin gray coating. These films contained oxides and hydroxides of Co and Sm (FIGS. 34(b) & (c)). On the other hand, oxides/hydroxides were not found in the metallic deposit obtained from the bath 1 (with 0.15M glycine) at 60° C. (650 mA/cm2) (FIG. 34(a)) and only a weak (11.0) peak was observed at 25° C. (50 mA/cm2) (FIG. 34(b)). The addition of glycine apparently suppressed formation of oxides and hydroxides. In the absence of glycine, bath 6 became unstable and white precipitates formed after 24 hours. Addition of glycine apparently stabilized the solution preventing the formation of hydroxides in the solution.

It is interesting that the Sm oxides included SmO in the deposits obtained from solutions without glycine (FIGS. 34(b) & (c)) indicating that Sm(II) may form during electrodeposition. Reduction potential of Sm2+ to Sm is much more negative, E°=−2.67 VSHE, than Sm3+/Sm2+ (E°=−1.55VSHE).W. M. Latimer, The Oxidation States of the Elements and Their Potentials in Aqueous Solution, Prentice-Hall, N.Y., pp. 286-295 (1952) Glycine forms complexes with Co and Sm ions enabling co-deposition of Co—Sm alloys. By addition of 0.15M glycine (bath 1), a relatively high Sm deposit content of 25 at % was obtained at 650 mA/cm2 and 60° C.

By addition of excess glycine (3M glycine, 60° C. bath 7), CDmax decreased (FIG. 32(f)) and Sm deposit content decreased (FIG. 33). At 60° C. and 400 mA/cm2, for example, the Sm content drops substantially from 14.7 to 2.9 at % as the concentration of glycine increased from 0.15 to 3M. Excess glycine may complex virtually all of the Co2+ and additional Sm3+ resulting in the formation of mononuclear complexes, Co(gly)3 and Sm(gly)3, at the expense of forming the heterodinuclear complexes required for the codeposition of Co—Sm alloys as proposed previously.[N. V. Myung, M. Schwartz, and K. Nobe, in Fundamental Aspects of Electrochemical Deposition and Dissolution, M. Matlosz, D. Landolt, R. Aogaki, Y. Sato, and J. B. Talbot, Editors, PV 99-33, p. 263, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999).]

(12) Effect of Ammonium Sulfamate Concentration

Ammonium sulfamate as supporting electrolyte in the plating bath has been used in a previous study [M. Schwartz, F. He, N. Myung, and K. Nobe, in Magnetic Materials, Processes, and Devices V. Applications to Storage and Microelectromechanical Systems (MEMS), L. T. Romankiw, S. Krongelb, and C. H. Ahn, Editors, PV 98029, p. 646, The Electrochemical Society Proceedings Series, Pennington, N.J. (1999)]. However, the effect with/without ammonium sulfamate on deposit properties (especially Sm content) has not been carefully studied yet. Bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine) and bath 8 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, 1M NH4 sulfamate) were used to study the effect of ammonium sulfamate as shown in Table 11. The Hull cell patterns of these deposits are given in FIG. 35.

TABLE 11
The effect of ammonium sulfamate
[Sm(NH2SO3)3][CoSO4][Glycine][NH4(NH2SO3]TCDmax
EXP #Bath(M)(M)(M)(M)(° C.)(mA/cm2)
311.000.050.1502550
660750
348125240
4060900
*Total charge = 100 C, applied current = 4.5 A, substrate area = 15 cm2 (10 cm × 1.5 cm), no agitation.

The presence of 1M ammonium sulfamate resulted in increased CDmax at 25 and 60° C. For example, at 25° C., the CDmax increased from 50 to 240 mA/cm2 by addition of 1M ammonium sulfamate (FIGS. 35(a) and (c)), but Sm content decreased, especially at 60° C. and higher CDs (FIG. 36). At 60° C., the highest Sm deposit content dramatically dropped from 25 to 12 at % (at 650 mA/cm2) by addition of 1M ammonium sulfamate.

(13) Results and Discussion of Pulse Current (PC) Electrodeposition

Effect of peak current density (PCD) and solution temperature: Since metallic deposits obtained at higher CD have higher Sm contents, PC electrodeposition, which enables higher PCD than DC, was used to increase the Sm deposit content. Bath 1 was used to study the effects of PCD and solution temperature on PC electrodeposition. Experimental conditions are given in Table 12, and Hull cell patterns of deposits are shown in FIG. 37.

TABLE 12
The Effect of Peak Current Density
TonToffFrequencyPCDmax.(or CDmax)
EXP #Bath(ms)(ms)Duty Cycle(Hz)T (° C.)(mA/cm2)
52110900.11025200
490.050.452k1160
56109010601000
530.050.452k2400
31DC2550
660750
*In PC, total charge = 50 C, substrate area = 7.5 cm2 (10 cm × 0.75 cm); at 25° C., applied current = 4.5 A, at 60° C., applied current = 7 A. In DC, total charge = 100 C, substrate area = 15 cm2 (10 cm × 1.5 cm), applied current = 4.5 A. Both DC and PC were no agitation

It was observed that PC electrodeposition has a larger maximum peak current density (PCDmax) than DC (FIG. 37); for example, at 25° C., the CDmax for DC was 50 mA/cm2 and the PCDmax for Ton=0.05 ms (duty cycle=0.1) was 1160 mA/cm2. As DC, the PCDmax increased with increased solution temperature. For instance, for deposits obtained from bath 1 at Ton=0.05 ms (duty cycle=0.1), PCDmax increased from 1160 to 2400 mA/cm2 by raising the solution temperature from 25 to 60° C.

At 25° C., because PC was able to obtain metallic deposits at higher PCD than DC, the highest Sm deposit content achieved by PC was greater than DC (FIG. 38). PC electrodeposition (Ton=0.05 ms, duty cycle=0.1) resulted in a maximum Sm content of 19 at % at 580 mA/cm2 and 25° C. At the same temperature (25° C.) DC electrodeposition had a maximum Sm deposit content of 5 at % at 25 mA/cm2.

On the other hand, at higher solution temperature (60° C.) PC did not result in a higher maximum Sm deposit content than DC electrodeposition; the maximum Sm deposit content by PC (γ=0.1, Ton=0.05 ms) was 9.5 at % (1900 mA/cm2) which was less than by DC (Sm=18.5 at % at 600 mA/cm2). At 60° C., although PC had higher PCDmax, the increase rate in Sm content for PCD (or CD)

SMcontentPCD(orCD),

was much lower for PC than for DC leading to this result. To sum up, at 25° C., a higher maximum Sm deposit content was obtained by PC than by DC electrodeposition; at 60° C., DC had a higher maximum Sm deposit content than PC.

Effect of Ton: Ton, relates to the total charge passed per pulse cycle; the longer Ton, the more charge passed for electrochemical reactions. All the deposits were obtained from 25° C. bath 1; the experimental conditions are given in Table 13; and the Hull cell patterns of deposits are shown in FIG. 39.

TABLE 13
The effect of Ton
ToffDutyFrequencyTPCDmax
EXP #BathTon (ms)(ms)Cycle(Hz)(° C.)(mA/cm2)
4910.050.450.12k251100
500.10.91k1100
5119100380
521090 10200
*Total charge = 50 C, substrate area = 7.5 cm2 (10 cm × 0.75 cm), no agitation; applied peak current = 4.5 A.

At shorter Ton higher PCDmax were obtained (FIG. 39). Shorter Ton resulted in lower Sm deposit content but since the metallic region was significantly extended (greater PCDmax), higher Sm content could be obtained. Deposits obtained from 25° C. bath 1 for Ton=0.1 ms, duty cycle=0.1, for instance, had 22 at % Sm content at 650 mA/cm2 (FIG. 40).

(14) Effect of Duty Cycle

Bath 1 was used to study the effect of duty cycle on electrodeposited Co—Sm alloys; experimental conditions are given in Table 14 and FIG. 41 show their Hull cell patterns.

TABLE 14
The effect of duty cycle
TonToffDutyFrequencyTPCDmax
EXP #Bath(ms)(ms)Cycle(Hz)(° C.)(mA/cm2)
11210.11.900.050500251300
1111.230.0757501250
500.90.11k1100
830.40.22k160
840.230.33k80
301DC50
(continuous)
*Total charge = 50 C, substrate area = 7.5 cm2 (10 cm × 0.75 cm), no agitation and applied peak current = 4.5 A.

Smaller duty cycle led to higher PCDmax For example, PCDmax dramatically increased from 50 to 1300 mA/cm2 by decreasing duty cycle from 1 to 0.05 (FIG. 41); Sm deposit content decreased (FIG. 42). For deposits obtained at 600 mA/cm2, Sm deposit content decreased from 21 to 7.5 at % by decreasing the duty cycle from 0.1 to 0.05.

3. Parametric Aqueous Electrodeposition Studies of Co—Sm Alloys

Iron group (IG)-rare earth (RE) alloys are known for their ferromagnetic and energy storage applications and resistance to aggressive environments. Co—Sm alloys, such SmCo5 and Sm2Co17, have already been commercialized for their high performance magnetic properties. These films have been prepared by sputtering [H. C. Theuerer, E. A. Nesbitt and D. D. Bacon, J. Appl. Phys., 40, 2994 (1969).], evaporation [V. Geiss, E. Kneller and A. Nest, Appl. Phys., A27, 79 (1982).], and plasma spraying [K. Kumar, D. Das and E. Wettstein, J. Appl. Phys., 49, 2052 (1978).].

SmCo5 alloys have very large coercivities as a result of its considerable magnetic anisotropy constant (ku) of about 107J/m3, and they also have high Curie temperatures. The latter enables high operating temperatures for permanent magnet applications as in magnetic coupling, sensors, nano and micro systems, servo motors, etc. Although CoSm alloys are expensive, their superior high temperature magnetic performance and reliability outweigh the costs for military and aeronautical/aerospace applications [M. Rassignol and J. P. Yonnet, Magnetism, II-Materials and Applications, E. T. de Lacheisserie, D. Gignoux and M. Schlenker, Eds., Chap. 15, Kluwer Academic Publishers, The Netherlands (2002).]. The development of an aqueous electrodeposition process for Co—Sm alloys will substantially lower manufacturing costs.

The high perpendicular coercivities of Co—Sm alloys make them eminently suitable for ultra high density information storage as in hard disk drives [M. H. Kryder, Applied Magntism, R. Gerber, C. D. Wright and G. Asti, Eds., p. 39, Kluwer Academic Publishers, The Netherlands (1994); R. C. O'Handley, Modern Magnetic Materials, Chaps. 13 & 17, John Wiley & Sons Inc., New York (2000); J. Sayama, K. Mizutani, T. Asahi, J. Ariake, K. Ouchi, S. Matsumura and T. Osaka, J. Magn. Magn. Mater., 287, 239 (2005).]. Iwasaki and Nakamura had earlier proposed (1977) perpendicular magnetic recording [S. Iwasaki and Y. Nakamura, EEE Trans. Mag., MAG-13, 1272 (1977).], and it has now been commercially realized.

Like other rare earth metals, electrodeposition of metallic samarium from aqueous electrolytes has not been achieved. Similarly, refractory metals such as W, Mo and V also have not been electrodeposited from aqueous media. However, electrodeposition of these metals from nonaqueous media can be done [N. Usuzaka, H. Yamaguchi and T. Watanabe, Mater. Sci. Eng., A99, 105 (1988).]. Although pure metals of W, Mo and V have not been electrodeposited from aqueous media, electrodeposition of alloys of W and Mo with the iron group metals (Ni, Fe, Co) have been readily done for over 60 years [M. L. Holt and M. L. Nielsen, Trans. Electrochem. Soc., 82, 193 (1942); H. J. Seim and M. L. Holt, Ibid, 96, 205 (1949).], IG-V binary and ternary magnetic thin film alloys by electrodeposition from aqueous solutions have been reported recently [C. Arcos, M. Schwartz and K. Nobe, Plat. Surf Finish., 90 (6), 46 (2003).].

More recently, IG-RE alloy electrodeposition from aqueous media has been achieved by our group by the use of glycine and other aminocarboxylates as complexers [L. Chen, M. Schwartz and K. Nobe, Proc. Electrochem. Soc., PV96-19, 239 (1996); M. Schwartz, F. He, N. Myung and K. Nobe, Ibid., PV98-20, 646 (1999); N. Myung, M. Schwartz and K. Nobe, Ibid., PV99-33, 263 (1999); M. Schwartz, N. Myung and K. Nobe, J. Electrochem. Soc., 151 (7), C468 (2004). 17. H. S. Cho, IEEE Trans. Magn., 33 (5), 2890 (1 997).]. The metal ions and glycine are known to faun hetero-nuclear glycinato coordinated complexes in aqueous solutions.

a. Experimental

Unless otherwise noted, the plating solutions consisted of 1M samarium sulfamate, 0.05M cobalt sulfate, 0.15M glycine as the complexer and 1M ammonium sulfamate as the conducting salt; also the total charge passed was 50 coulombs. The pH value of the plating bath was 5.2 as measured at 25° C. Various current densities and bath temperatures (25-60° C.) were used to obtain deposits; solutions were not agitated during electrodeposition.

An EG&G PAR potentiostat, model 173, served as the power source for electrodeposition. Brass panels (2×1.9 cm) served as cathodes and a platinum sheet (3×6 cm) was used as the anode. The brass panels were mechanically cleaned, soaked in alkaline solution for 10 min., rinsed in deionized water and immersed in 10% HCl for 30 seconds.

The ratios of the samarium and cobalt deposit content was determined by energy dispersive x-ray analysis (EDX); deposit content of cobalt was also determined by atomic absorption spectrophotometry (AA). The microstructure, crystal orientation and grain size were determined by x-ray diffraction (XRD) and surface morphology by scanning electron microscopy (SEM). Magnetic properties were determined by a vibrating sample magnetometer (VSM), model 1660 ADE Tech.).

b. Results and Discussions

The experimental data presented in the figures are for deposits which had metallic appearance. In each case, at higher current densities (CD), deposits were non-metallic and these results are not shown.

FIG. 43 shows the effect of increased solution temperature on the applied CD range. At 25° C., the effective CD for metallic deposits was limited to 350 mA/cm2, whereas for a solution temperature of 60° C., the CD could be increased to 700 mA/cm2, resulting in increasing deposit Sm content from 8 to 17 at %. Although the (linear) deposition rate was greater at 25° C., extrapolation indicated that at 450 mA/cm2 the deposition rate at 60° C. would have exceeded that of the rate at 25° C. Elevated solution temperatures permit higher CDs depositing metallic appearing deposits with deposit Sm content reaching 17.2 at % at CD˜700 mA/cm2 and a temperature of 60° C. The current efficiencies (CE) drop sharply between 50 mA/cm2 (20%) and 200 mA/cm2 (6%) with little change between 200 mA/cm2 and 350 mA/cm2 in the 25° C. solution, whereas the CE decreases almost linearly in the 60° C. solution with an apparent slope of about 1%/100 mA/cm2 (FIG. 44).

With increasing solution temperature from 25-60° C., the deposit Sm content decreases almost linearly for CDs of 100 mA/cm2 and 300 mA/cm2, the latter being consistently 4% higher (FIG. 45). However, the CEs were higher for the lower applied CD (FIG. 46).

Additional parametric studies were performed to assess the effects of agitation, the solution concentration of Sm sulfamate and glycine, and the presence of NH4 sulfamate on the deposit Sm content. Agitation had a greater effect on the deposit Sm content at higher CDs while it generally increased with decrease in solution concentration of the Sm salt from 1M. An increase in the glycine concentration from 0.15M decreased the deposit Sm content. Addition of the conducting salt (NH4 sulfamate) decreased the deposit Sm content.

Magnetic saturation (Ms) values of the electrodeposits are quite close but slightly lower than those of sputtered Co—Sm thin films at equivalent Sm contents (FIG. 47). This is indicative of the metallic nature of the electrodeposited Co—Sm alloys. The electrodeposition results in this figure were obtained from solutions with no conducting salt (NH4sulfamate). The closed circle points represent electrodeposits obtained at 60° C. with 50 coulombs of charge passed. The open points represent thicker deposits (500 coul. charge passed). Higher Sm content could be obtained for thicker deposits from solutions in the absence of NH4sulfamate (˜33 at. % Sm, open diamonds) than in its presence (23 at. %Sm, closed circles). Pulse current electrodeposition (open triangles) produced metallic deposits (20 at. %Sm) for a duty cycle (7) of 0.1. Magnetic saturation of non-metallic deposits were substantially lower than metallic deposits of the equivalent Sm content. For example, for 17 at %Sm, Ms of the metallic deposit was 4 times higher than the non-metallic deposit. FIG. 48 shows that the deposit coercivities (Hc) in the parallel direction increase only slightly with increase in deposit Sm content, i.e., with increased CD, and is not affected by solution temperature. However, deposit coercivities decrease sharply in the perpendicular direction with increasing CD in deposits from 25° C. solutions, but a linear decrease with a negative slope from 60° C. solutions. Perpendicular coercivities are significantly higher than in the parallel direction. Thus, heat treatment of Co—Sm deposits leads to substantially higher coercivities.

The deposit topography is affected by the applied CD and probably also solution temperature. There is increased surface roughness as a result of increased CD. FIGS. 49a,b show that the absence of the conducting salt (NH4sulfamate) dramatically results in a much smoother surface than in its presence. X-ray diffraction spectra (XRD) (FIG. 50) indicate the 2.1 at. % Sm deposit (100 mA/cm2) appears amorphous and the 10.4 at. % Sm deposit (500 mA/cm2) exhibits crystalline structures with (200) phase of Sm2Co17 alloy composition. In the absence of NH4 sulfamate in the bath, the (201) phase of the SmCo5 alloy as well as the Sm2Co17 (200) phase appear in the deposits.

c. Conculsions

Sm content of metallic deposits of Co—Sm can be increased at higher CDs from higher temperature solutions. Furthermore, significantly higher deposit Sm content can be obtained from solutions in the absence than in the presence of the conducting salt (NH4 sulfamate). Co—Sm deposits with 33 at % Sm have been obtained at 500 mA/cm2 and a solution temperature of 60° C. Magnetic saturation of electrodeposits were close to those of sputtered deposits. Perpendicular coercivities were substantially higher than parallel coercivities for Co—Sm electrodeposits. Heat treatment of deposits should result in an order of magnitude increase in perpendicular coercivities.

4. Coordination Chemistry in the Electrodeposition of IG-V, W and Mo Alloys from Aqueous Carboxylate Solutions

Polycarboxylates, hydroxycarboxylates and aminocarboxylates are well known additives functioning as complexing agents for the electrodeposition of single metals and alloys from aqueous plating baths. Tartrates and citrates are extensively employed in electrodeposition of alloys, including the deposition of alloys containing elements such as the refractory metals: W, Mo, V. Although these individual metals cannot be electrodeposited from aqueous media, alloys with the iron group metals (IG) have been electrodeposited from aqueous solutions [A. Brenner, Electrodeposition of Alloys, Vol. II,. Acad. Press (1963).

Brenner et al. and Holt and his co-workers have studied the electrodeposition of IG-W and —Mo alloys from aqueous solutions [A. Brenner, P. Burkheard and E. Seegmiller, J. Iles. NBS, 94, 351 (1947); L. E. Vaaler and M. L. Holt, Trans. Electrochem. Soc. 90, 43 (1946); L. E. Vaaler and M. L. Holt, Ibid., 94, 50 (1948); W. E. Clark and M. L. Holt, Ibid., 94, 244 (1948); M. H. Lietzke and M. L. Holt, Ibid., 94, 252 (1948); R. F. McElwee and M. L. Holt. I Electrochem. Soc., 99 (2), 48 (1952).]. More recent work on the electrodeposition of IG-X alloys has been reported for electroplating baths containing citrates as complexers, with some consideration given to the nature and structure of the organometallic complexes involved [M. Schwartz, C. Arcos and K. Nobe, Plat. Surf. Fin., 90 (6) 46 (2003); E. J. Podlaha and D. Landolt, J. Electrochem. Soc., 143, 885 (1996); Ibid., 143 (11) 893 (1996); Ibid., 144 (5) 1672 (1997); O. Younes and E. Gileadi, Ibid., 149, C100 (2002); O. Younes-Metzler, L. Zhu and E. Gileadi, Electrochim. Acta, 48, 2551 (2003).]. There has been a growing interest in determining the structures of these refractory metal-hydroxycarboxylato—complexes by spectroscopic experiments for biological, physiological applications [M. Tsaramyrsi, M. K. aliva, A. Salifoglou, C. P. Raptopoulou, A. Terzis, V. Tangoulis and J. Giapintzakis, Inorg. Chem., 40 (23), 5773 (2001); T. Kiss, P. Buglyo'. D. Sanna, G. Micera, P. Decock and D. Dewaele, Inorg. Chem. Acta, 239, 145 (1995); 9. Z-H Thou, H-L Wan and K-R Tsai, 1 Chem. Soc., Dalton Trans., 4289 (1999). Z-H Thou, H-L Wan and K-R Tsai, Inorg. Chem., 39, 59 (2000); Z-H Thou, S-Y Hou and H-L Wan, J. Chem. Soc., Dalton Trans., 1393 (2004).].

Since the early 1990s, this laboratory has been investigating the electrodeposition of IG-V alloys in order to improve the physical properties as well as the corrosion resistance of the high magnetic moment 90Co10Fe alloys, as suggested by Liao [S. H. Liao, IEEE Tran. Magn 23, 2981 (1987).]. In addition, the electrodeposition of V-Petinendur (49Co49Fe2V) was investigated because of its excellent magnetic properties; it has not had wide commercial applications because of its high manufacturing costs, however [G. Y. Chin and J. H. Wernick in Ferromagnetic Materials, vol. 2, p. 168, E. P. Wohlfarth, Ed., North-Holland, Amsterdam, (1980).]. Development of a commercial electrodeposition process would sharply reduce these costs and greatly expand applications in miniaturized electronic devices.

Some of our interest in the electrodeposition of the IG-X alloys (X=V. W and Mo) dates to the early work (1948) of one of us (M. S.) who developed a commercial Co—W plating process using an ammoniacal citrate bath. He found that when the Co and WO42− salt solutions are mixed, a cobalt tungstate precipitate forms which dissolves with the addition of citrate. His experimental results lead him to conjecture that both Co(II) and W(VI) are coordinated in the same complex with deprotonation of the carboxylate forming a heteronuclear biscitrato—complex. Subsequently, later work by Zhuravleva and co-workers on the complexation of oxyvanadium and IG ions with citric acid indicated the existence of homo-dinuclearVO-biscitrato- and hetero-dinuclear IGbiscitrato- complexes [Y. I. Sal'nikov, F. V. Devyatov, N. E. Zhuravleva, D. V. Golodnitskaya, Jh. Neorgan Khimi., 29, 2273 (1984); A. N. Glebov, Y. I. Sal'nikov, N. E. Zhuravleva, V. V. Chevela and P. A. Vasil'ev, Ibid., 27, 2146 (1982); N. E. Zhuravleva and Y. I. Sal'nikov, Proc. Tech. Inst. Chem. and Chem Tech., 32 (2) 25 (1989).].

Plating baths are presented in Table 15, Table 16, Table 17, and Table 18. NH3 (aq) and HCl (or sulfamic acid) were used to adjust solution pH. Plating baths were maintained at room temperature and were quiescent or mechanically stirred, as required. DC was provided by a PAR potentiostat or a Kraft Dynatronix power supply (Model DPR 20-5-10). Total charge passed was controlled to produce desired film thicknesses.

Brass panels served as cathodes and the anode was a cobalt (or Pt) sheet. The brass panels were scrubbed with Alconox, rinsed in deionized water and then activated in concentrated HCl. The deposits were dissolved in concentrated HNO3 for analysis. Cobalt and vanadium deposit contents were determined by atomic absorption spectrophotometry (AA). Cobalt was also determined by energy dispersive spectroscopy (EDX). Tungsten and molybdenum were analyzed by gravimetric methods as well as EDX. Scanning electron microscopy (SEM) and x-ray diffraction were used to characterize the deposits. Magnetic properties of the electrodeposits were obtained with vibrating sample magnetometer (VSM) (ADETech, Model 1660).

a. IG-Rare Earth (SM) Coordination Compounds

Initial investigation of the electrodeposition of IG-RE alloys was reported in 1996 [L. Chen, M. Schwartz and K. Nobe, in Electrodeposited Thin Films. M. Paunovic and D. A. Scherson, Editors, PV96-19, p. 239, The Electrochem. Soc. Proceedings Series, Pennington, N.J. (1996).]. Hydroxycarboxylate and other carboxylate salts as complexers appeared to be inferior to glycine and other amino acid salts. The bipolarity, i.e. zwitterionic properties, of glycine (and other aminoacids) results in ionic complexes at the cathode surface as the pH fluctuates. As a result, structures and deposition mechanisms of IG-RE dinuclear diglycinato- and triglycinato-coordination complexes have been considered as the vehicles for the co-deposition of the IG-RE alloys [M. Schwartz, M. V. Myung and K. Nobe, J. Electrochem. Soc., 151. C468 (2004).]. Continued development indicates the deposit Sm contents and resulting magnetic properties could be varied, depending on the solution composition, CD and temperature [J. C. Wei, M. Schwartz and K. Nobe, Trans. Electrochem. Soc. (in press).]. Table 15 summarizes typical solution composition ranges and with the deposition conditions selected, result in wide variations in deposit Sm content. The glycine: Co ratio in these solutions was 3:1 with Sm3+ concentration in excess. Solutions with lower Co and Sm concentrations resulted in higher deposit Sm contents (15-18 a/o) than more concentrated solutions. In the latter solution, FIG. 51 shows the deposit Sm content increased linearly with increasing CD while the CE decreased sharply as CD increased to ˜100 mA/cm2, reaching a plateau with increased CD, indicating a possible relation of hydrogen evolution to deposit content.

TABLE 15
Representative plating baths for Co—Sm alloys
Co2+ (as sulfamate)0.06M-0.12M
Sm3+ (as sulfamate) 0.3M-0.9M
Glycine0.18M-0.36M
NH4(NH2SO3)1.0M
NH4OHpH 6.5-7.0

TABLE 16
Representative plating baths for IG-V alloys
BinaryTernary
Co2+0.3M0.15-0.3 M
VO2+0.15-0.30.15-0.17 M
Fe2+0.03M-0.15M
(Ni2+)(0.05M)
Na3C6H5O70.25-0.35 M0.25 M
H3BO30.1 M
NH4Cl1.0 M
NH4OHpH 6.0-7.5

b. IG-Vanadium Coordination Compounds

Electrodeposition of binary and ternary IG-V alloys from citrate solutions has been reported [M. Schwartz, C. Arcos and K. Nobe, Plat. Surf. Fin., 90 (6) 46 (2003); B. Y. Yoo, dissertation, UCLA, 2003; also unpublished data, UCLA. 2004.]. Although “representative” solution compositions are given in Table 16, wide variations in concentrations provide flexibility in the resulting deposit compositions and their magnetic properties. In binary alloy deposits, the V contents decreased: Co>Fe>Ni and the magnetic properties varied with the deposit V content; magnetic saturation decreased and coercivities increased with increasing deposit V content (Table 19). Thus, the ability to control deposit composition and their magnetic properties provides “tailor-made” deposits for various electronic applications. FIG. 52 indicates the wide variations in Co—V electrodeposits vs CD with deposit Co content increasing with increasing CD, whereas Fe-V electrodeposits exhibit very little change in deposit compositions vs CD, with deposit V contents <2 w/o over the CD range, 5-50 mA/cm2 [B. Y. Yoo, dissertation, UCLA, 2003; also unpublished data, UCLA. 2004.]. Magnetic saturation, Bs, of ternary CoFeV electrodeposits exceeded that of the binary CoV and FeV electrodeposits. Ternary 48Co49.8Fe2.2V electrodeposits approximating bulk 2V-Permendur (48Co48Fe2V) exhibited a magnetic saturation Bs of 2.32 T. Liao indicated a 90Co 10Fe electrodeposit had a magnetic saturation of 1.9 T, exceeding that of Permalloy (80Ni20Fe), Bs=1.0 T; however, corrosion resistance of the alloy was inferior [S. H. Liao, IEEE Tran. Magn 23, 2981 (1987).]. As shown in Table 19, addition of V to the binary electrodeposit (87.3Co8.6Fe4.1V) improved corrosion resistance and increased magnetic saturation, Bs=2.20 T [B. Y. Yoo, dissertation, UCLA, 2003; also unpublished data, UCLA. 2004.].

TABLE 17
Representative plating baths for Co—W alloys
Brenner
Holt (3)(2)Brenner (2)MS (12)Gileadi (6)
[Co2+] (M)0.2130.60.420.12*(Ni)
[WO42+] (M)0.2130.30.130.12*
[Cit3−] (M)0.3141.20.48*
[Tart2−] (M)1.42
NH4Cl0.50.940.5 *
NH4OH7-8.58.58.58.5-9   8
(pH)
% W, a/o>17.57.77.7-9.6 7.9-17.512-15
% W, w/o>40~2020-5020-4030-35
T (° C.)70≧40≧9065-8025

TABLE 18
Representative plating baths for IG-Mo alloys
Holt (22)Landolt (5)
IG [Fe3+], [Co2+], [Ni2+] (M)0.30.2 (Ni)
Na2MoO4 (M)0.02-0.0750.005-0.05 
Na3C6H5O70.30.25-0.95
NH4OH (pH)10.5 9.7-10.2
T (° C.)2525-40

TABLE 19
Magnetic properties of selected electrodeposited IG-V, W, Mo alloys
CoBs
(wt %)Fe (wt %)V (wt %)W (wt %)Hc (Oe)(T)Reference
9191181.43(20)
97.82.2251.93(20)
6733571.28(19)
87.38.64.1582.20(20)
48.049.82.2462.32(20)

Nikolova and Nikolov suggested mononuclear and dinuclear oxyvanadium citrato complexes. Based on indirect evidence, potentiometric experiments and IR spectra, they indicated the protonated hydroxo-group may be involved in 1:1 mononuclear citratocomplex (stability constant=6.6×108) and a dinuclear citrato- complex with the hydroxy group or carboxylato group bridging the two VO's (stability constant=3.2×1011) [B. M. Nikolova and G. St. Nikolov,./.Inorg. Nucl. Chem., 29, 1013 (1967).].

Zhuravleva and co-workers studied dinuclear IG biscitrato-complexes and heteronuclear IG-(VO)2-biscitrato-complexes [Y. I. Sal'nikov, F. V. Devyatov, N. E. Zhuravleva, D. V. Golodnitskaya, Jh. Neorgan Khimi., 29, 2273 (1984); A. N. Glebov, Y. I. Sal'nikov, N. E. Zhuravleva, V. V. Chevela and P. A. Vasil'ev, Ibid., 27, 2146 (1982); N.E. Zhuravleva and Y. I. Sal'nikov, Proc. Tech. Inst. Chem. and Chem Tech., 32 (2) 25 (1989).]. Salni'kov et al. investigated mononuclear, dihomonuclear and heteronuclear biscitrato-complexes of Ni and Co, the latter being partially deprotonated or completely deprotonated [Y. I. Sal'nikov, F. V. Devyatov, N. E. Zhuravleva, D. V. Golodnitskaya, Jh. Neorgan Khimi., 29, 2273 (1984).]. Glebov et al. indicated dinuclear (VO)2 complexes only with hydroxocarboxylic acid, e.g., citric, tartaric and malic acids [A. N. Glebov, Y. I. Sal'nikov, N. E. Zhuravleva, V. V. Chevela and P. A. Vasil'ev, Ibid., 27, 2146 (1982).]. The depicted dinuclear biscitrato-coordination compound indicates protonated hydroxy group oxygens bridging the two V atoms. Zhuravleva and Salni'kov obtained a heteronuclear biscitrato complex by reacting individual Cu citrate and VO citrate solutions [N. E. Zhuravleva and Y. I. Sal'nikov, Proc. Tech. Inst. Chem. and Chem Tech., 32 (2) 25 (1989).].

c. IG-W Coordination Compounds

Holt and his students investigated the electrodeposition of IG-W alloys from ammoniacal citrate solutions [L. E. Vaaler and M. L. Holt, Trans. Electrochem. Soc. 90, 43 (1946); L. E. Vaaler and M. L. Holt, Ibid., 94, 50 (1948); W. E. Clark and M. L. Holt, Ibid., 94, 244 (1948); M. H. Lietzke and M. L. Holt, Ibid., 94, 252 (1948); R. F. McElwee and M. L. Holt. I Electrochem. Soc., 99 (2), 48 (1952).]. After complexation, atomic hydrogen reduction has been suggested to explain the co-deposition of IG-W alloys by proposing a two-step reduction hypothesis involving alternating deposition of the IG species which catalyzed the reduction of the tungstate ion and resulted in a laminar deposit, based on polarographic studies and cathode potential measurements, the latter being lower in solutions containing tungstate ions.

Younes and Gileadi concluded a heteronuclear Ni-W monocitrato complex with the carboxylate triply deprotonated is the precursor for the electrodeposition of the Ni—W alloy [O. Younes and E. Gileadi, Ibid., 149, C100 (2002); O. Younes-Metzler, L. Zhu and E. Gileadi, Electrochim. Acta, 48, 2551 (2003).]. The complex is the result of the reaction of individual Ni and W citrate complexes. A similar reaction with a Ni-biscitrato-complex was considered unlikely because both reacting complexes would be highly charged. They also reported that elimination of NH4 salts and NH3 (aq.) from similar solutions resulted in amorphous deposits with increased deposit W contents, but substantial CE reduction.

Brenner et al. indicated an inorganic Co—W solution (no citrate) resulted in deposits containing 20-27 wt.% W at CDs 20-50 mA/cm2, respectively [A. Brenner, P. Burkheard and E. Seegmiller, J. Iles. NBS, 94, 351 (1947).]. Without the presence of NH4 salts, the solubility of Co and W is reduced and the CE is quite low, making the solution unsuitable for practical applications. In the tartrate complexed Co—W solution (Table 17), the presence of alkali cations (Na, K tartrate) resulted in diminished deposit W contents as compared to ammonium tartrate solutions, another indication of the positive effect of the presence of NH4 ions.

d. IG-Molybdenum Coordination Compounds.

Holt and students extended their studies of the electrodeposition of IG-W alloys to IG-Mo alloys from equivalent solution compositions, as shown in (Table 18), [J. Seim and M. L. Holt, Trans. Electrochem. Soc., 95, 205 (1949); D. W. Ernst, R. F. Amlie and M. L. Holt, J. Electrochem. Soc., 102 (8), 461 (1955); D. W. Ernst and M. L. Holt, Ibid., 105 (11), 686 (1958).]. Initially, Hull Cells were utilized to determine pH, temperature and CD ranges for bright, metallic deposits and resulting deposit composition and CEs. Various ligands were classified as “good” or “poor”: For Ni—Mo, citrate and tartrate were considered “good”; for Co—Mo, the “good” ligands were extended to include malate (and malic acid) and glycolic acid. Sodium citrate was considered superior to citric acid or ammonium citrate.

Similar to IG-W electrodeposition, deposit Mo contents varied with the co-deposited IG metal: Fe (59%)>Co (40%)>Ni (20%). The CEs were inversely related to the deposit Mo contents: Ni (75-85%)>Co (50-60%)>Fe (10-20%), which seems to indicate a role for adsorbed hydrogen atoms in the deposition process.

Podlaha and Landolt studied the effect of varying the solution Ni:Mo ratios on mass transport and kinetically controlled processes with rotating cylindrical electrodes [E. J. Podlaha and D. Landolt, J. Electrochem. Soc., 143, 885 (1996); Ibid., 143 (11) 893 (1996); Ibid., 144 (5) 1672 (1997).]. In addition to the metal salts, the ammoniacal solution contained sodium citrate; no conducting salts were added (Table 18). The concentration of NH3 (aq.) appears to be critical with respect to the deposit Mo content. “High” NH3 (aq.) concentrations resulted in high CEs 90% and lower deposit Mo content; most of the data presented were obtained with 0.28 M NH3 (aq.). A deposition mechanism whereby Mo is deposited from an intermediate (Ni—Mo) complex either adsorbed or dissolved in the solution is proposed with Ni depositing independently and considered the catalyst for the electrodeposition of the alloy. The model is extended to include codeposition of Co—Mo and Fe—Mo and explain the differing effects of the IG species on the deposit Mo content by including the additional presence of a single IG complex with a two-step reduction.

Zhou and collaborators prepared and determined the structures of various Mo, W, and V citrato-coordination compounds, which are precursors to FeMo cofactors in nitrogenase protein catalysts, using IR and NMR spectroscopies and Xray diffraction [9. Z-H Thou, H-L Wan and K-R Tsai, 1 Chem. Soc., Dalton Trans., 4289 (1999). Z-H Thou, H-L Wan and K-R Tsai, Inorg. Chem., 39, 59 (2000); Z-H Thou, S-Y Hou and H-L Wan, J. Chem. Soc., Dalton Trans., 1393 (2004).]. Tridentate Binuclear biscitrato-complexes convert to bidentate mononuclear biscitrato ions with excess citric acid (pH 3.5). Each ligand is coordinated as a bidentate ligand through the deprotonated central carboxylato- and vicinal hydroxy-groups with the terminal carboxylate groups uncomplexed. The structure of the tridentate citratocomplexes depends on the degree of protonation of the reacting ligand and solution pH. Mononuclear monocitrato-complex converts to dinuclear biscitrato-complex with the two metal atoms connected via an oxygen bridge and two oxygen atoms covalently bonded to each metal atom in lower pH solution with valencies depending on whether or not the terminal carboxylate groups are protonated.

e. Structure of IG-RE-Glycine Complexes

The zwitterionic structure of glycine (gly) protonates or deprotonates depending on the solution pH, as shown FIG. 53a. Glycine usually coordinates through the carboxylate groups forming O-M-O bonds; it can also form O-M-N and N-M-N bonds [C. P. Sinha, Complexes of the Rare Earths, Pergamon Press, 1966.]. Single metal (i.e., IG2) complexes are probably mononuclear monoglycine bidentate chelate complexes. For homo- or hetero-dinuclear complexation, multiligand complexes are required through N-MI-N and O-M2-O bonds. Both IG and RE cations are likely in a hetero-dinuclear trisglycine complex as in FIG. 53b. However, because of the bipolar character of glycine, it can simulate polymerization by electrostatic attraction between deprotonated zwitterions, mimicing dip eptide or trip eptide structures through O—N “bonds”. In the case of the heteronuclear IG-RE complex, quasi-digly (FIG. 53c) or trigly (FIG. 53d) complexes can be formed and adsorbed at the cathode surface as the cathode film pH fluctuates.

f. Structure of IG-VO-Biscitrate Complexes

Citric acid can form di-, tri- or tetra-dentate complexes with metal ions, depending on the degree of deprotonation. The structure of the binary IG-VO citrate complex (FIG. 54a) has one carboxylate group protonated with the hydroxy-hydrogen of each citrate ligand attracted to the VO, possibly forming H-bonds. The depicted complex structure shows partial deprotonation of the citrate moiety, resulting in a zero-valent complex. With continued deprotonation, the complex becomes anionic with increasing negative charges. Thus, the degree of protonation/deprotonation of the carboxylatogroups is dependent on solution pH. As the pH in the cathode film fluctuates slightly, the forms of the coordination complex may equilibrate. FIG. 55 a indicates the hydroxy hydrogen from each citrate ligand reacts with the oxyvanadium (VO) component in step 2, leaving the IG-V citrate complex [IGIIVIV (C6HSO7)2]0 to be step-wise reduced by hydrogen atoms to deposit the IG-V alloy.

FIG. 54b shows the structure of the ternary IG-IG-V-citrate complex. After removal of the oxygen in the oxyvanadium (VO) component, as the binary complex, a similar sequential reduction process occurs to electrodeposit the ternary alloy (FIG. 55b).

g. Structure of IG-W (Mo) Biscitrate Complexes

The tungstate (molybdate) ion reacts with citric acid and the Id′ ion to form the IG-W(Mo) citrate complex (FIG. 56). Similar hydroxy-H bonds and reacts with one ofthe (MvI00) oxygens converting to the complex [IGIIMVI O(C6HSO7)2]0. FIG. 57, as FIG. 55, shows the similar reduction of the IGII and MVI in the IG-M citrate complex to the IG-M alloy.

h. Co—W/CR Composite Coating

Although there is much interest today on product miniaturization, such as thin film magnetic devices, utilizing these electrodeposited IG-refractory metal alloys from complexed solutions, and the structures of V, W and Mo (poly)homonuclear and (poly)heteronuclear polycitrato-complexes for physiological processes, as discussed above, we also point out potential industrial applications for thicker electroplated alloys, based on physical, mechanical and high temperature properties such as corrosion and wear resistance and the ability to undergo hardening with thermal treatment. FIG. 58 illustrates the potential of producing unique composite electrodeposits for specific applications. In this example, corrosion resistance, wear resistance and high temperature properties were required [M. Schwartz in Handbook of Deposition Technologies for Films and Coatings, R. F. Bunshah, Ed., Chpt 10, Noyes Publ. (1994).]. The composite consists of a 80Co-20W deposit (56 μm) from the solution in Table 17 +Cr from CrO3/H2SO4 solution (30 μm)+Co—W (20 μm)+Cr (thin deposit to protect deposit edges). The specimens were subjected to high temperature environment in air (to determine oxidation resistance) and a carburizing atmosphere, attempting to diffuse C into the Co—W deposit to form and disperse carbide in the deposit. To obtain good adhesion of these layers, a thin cobalt strike (S) was deposited on the steel substrate and between each deposit to improve deposit adhesion.

5. Direct Current (DC) Electrodeposition Studies

DC electrodeposition using a parallel electrode system was investigated to determine, more precisely, the optimum operating and aqueous bath conditions, estimated by the Hull cell studies, to obtain high Sm content, metallic Co—Sm alloys. The purpose is to obtain stoichiometric compositions of Co—Sm alloys to form intermetallic Sm2Co17 (10.5 at %Sm) and SmCo5 (16.7 at %Sm), and, in addition, Sm2O7 (22.2 at %Sm) and SmCo3 (25 at %Sm) after appropriate heat treatment procedures.

DC electrodeposition of Co—Sm alloys from aqueous solutions has been studied at various operating conditions (e.g., current densities, temperatures, fluid dynamics and pHs) and plating baths (concentration of Sm sulfamate, glycine, Co sulfate, NH4 sulfamate supporting electrolyte). Morphologies, crystalline structures, preferred orientation and microstructures of deposits from different electrodeposition parameters were studied and correlated to their magnetic properties.

Samarium deposit content increased with increasing current density. Increased solution temperatures from 25 to 60° C. effectively extended the CD to obtain metallic deposits (CD)max from 50 to 500 A/cm2 leading to a high Sm deposit content of 32 at % from a bath consisting of 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine at 60° C. At low CD of 10 & 50 mA/cm2 and 25° C., increasing solution pH increased Sm deposit content and then reached a constant. At 60° C., Sm content increased with increased solution pH and at 100 and 300 mA/cm2, Sm content reached maximum at pH 5 and 4, respectively. Rotating disk electrode (RDE) results suggest mass transfer effects in the deposition of Co.

Decreased Sm sulfamate concentration increased Sm deposit content. However, high content of Sm hydroxide/oxide was found indicating that the increase of Sm content was due to formation of Sm hydroxide/oxide in deposits. Highest Sm contents were obtained at glycine concentration between 0.1M (glycine: Co2+=2:1) and 0.15M (glycine: Co2+=3:1). Addition of NH4 sulfamate resulted in decreased Sm content.

Crystal structures were dependent on Sm content. Deposits changed from crystalline to non-crystalline structures by increased Sm deposit content. Crystal structures of electrodeposited Co—Sm crystallites were dominated by hcp phases. Higher Sm content deposits were usually accompanied with more microcracks due to higher internal stress by lattice distortion. Lower microcrack densities were found in the deposit obtained at higher solution temperatures.

Magnetic properties of electrodeposited Co—Sm alloys are strongly dependent on alloy composition, crystal structure and particle size of deposits. Increased Sm content resulted in deposits changing from crystalline to non-crystalline structures and decreased grain size. Magnetic saturation (Ms) decreased linearly with increased Sm deposit content and was in agreement with sputtered films. On the other hand, deposits with high oxide/hydroxide content had much lower Ms values. For deposits obtained at different conditions, perpendicular coercivity varied in the range of 150 to 1160 Oe, and parallel coercivity fluctuated between 50 and 150 Oe. Hc decreased sharply but changed little in the in-plane direction by increase in Sm content. The significant change in Hc may be the result of the considerable decrease in particle size in the in-plane direction. Parallel squareness (Mr/Ms) were higher than perpendicular squareness indicating the preferred magnetization direction lies on the deposit plane. Squareness decreased with increased Sm content for deposits changing from crystalline to non-crystalline structures.

In Hull Cell studies, we have already learned about how electrodeposition parameters qualitatively affected Sm deposit content and CDmax (the maximum CD for metallic deposits region) by Hull cell study. In addition, proper plating solutions and operating conditions of obtaining high Sm content deposit (26 at % Sm) by DC electrodeposition have also been achieved. In DC electrodeposition studies, we not only made deposits by parallel electrodes to confirm the Hull cell results but also studied the impacts of these parameters on the properties of electrodeposited Co—Sm alloys. Alloy properties, such as composition, crystal structure, morphology, microstructure and magnetic properties (i.e. coercivity Hc, saturation magnetization Ms and squareness S), varied by DC electrodeposition parameters will be studied and correlated. These studies are very helpful for not only getting a better understanding about the dependence of alloy properties on electrodeposition parameters but also providing important information to develop the mechanism or the electrodeposition of Co-Sin alloys in the future.

To understand the magnetic behavior of the electrodeposited Co—Sm thin films, it is necessary to study the dependence of Co—Sm alloy composition on operating conditions (i.e., current density, solution temperature, solution pH, and fluid dynamics) and plating solutions (i.e., concentration of Sm sulfamate, glycine, NH4 sulfamate and types of supporting electrolytes) and the dependence of morphology (i.e. roughness, cracks and pitting), crystalline (i.e., crystal structure, orientation, grain size and microstructure) on these parameters, then correlate these properties to the magnetic behaviors of electrodeposited deposits.

At the beginning of each section, the effect of parameter on alloy composition and current efficiency will be quantitatively measured and analyzed. The amount of Co and Sm deposition and H2 gas evolution will be showed in normalized charge diagrams. The second part mainly focused on the effect of parameters on crystal structures, orientations, grain size, morphologies and microstructures of deposits by their XRD patterns and the SEM micrographs. Finally, the effect of parameters on deposit magnetic properties (i.e., Hc, Ms and S) obtained from hysteresis loops of deposits will be studied and correlated to the other alloy properties discussed in the first and the second part. Magnetic properties of CoSm alloys as deposited will be revealed in this study.

The goals of the parametric studies of DC electrodeposition were: Confirming Hull cell results and obtaining high Sm deposit content Co—Sm alloys; Studying the dependence of alloy properties on electrodeposition parameters; Revealing the dependence of magnetic properties of electrodeposited Co—Sm alloys on other alloy properties (i.e., alloy composition, crystal structures, grain size . . . )

FIG. 59 shows the experimental flowchart of a DC experiment which mainly includes four parts: pretreatment of cathode, DC electrodeposition, post-treatment of specimen and characterization. Setup of DC electrodeposition, rotating disk electrode, design of experiments, pretreatment and post-treatment and characterization and analysis of specimens will be described in the following discussion.

a. Setup of DC Electrodeposition

FIG. 60 shows the setup of a DC electrodeposition system. An EG&G PAR potentiostat, model 173, served as the power source for electrodeposition. A coulometer to measure the total applied charge passed during the electrodeposition. The deposits were obtained in a 250 mL beaker filled with the plating solution. Tapped brass panels with 3.8 cm2 (2×1.9 cm) deposit area served as cathodes and a platinum sheet (3×6 cm) was used as the anode; in this experiment, the distance between the cathode and the anode is 4 cm. A saturated calomel electrode (SCE) served as a reference electrode to measure the cathodic potential during the electrodeposition process. A shielding panel with a 2×2 cm opening window, designed by the simulation result of ANSYS (a finite element analysis software), was placed between the cathode and anode to minimize the thickness variation of deposit by uniformizing the current density distribution on the cathode.

b. Setup of Rotating Disk Electrode (RDE) System

The setup of a RDE system shows in FIG. 61. A brass disk (diameter of 6.4 mm) was inserted into the cavity at the bottom of the Teflon rod; a copper rod was inserted into that Teflon rod from the other side and all the way down to touch the back side of the brass disk to make electrical contact. The RDE was jointed with a stainless rod by a coupler, and the stainless rod was rotated by a motor through a transmission belt. The RDE was dipped into the electrolyte with the brass disk facing down to the Pt panel located at the bottom of the plating cell. The distance between cathode and anode is about 2 cm. A SCE served as a reference electrode to measure the cathodic potential during electrodepositions.

C. Design of Experiments

In the study of DC electrodeposition, operating conditions, such as current density, temperature, pH, fluid dynamics will be varied. Various current densities (2-500 mA/cm2), bath temperatures (25-60° C.), pH values of plating solutions (2-6), fluid dynamics (RDE. 0-4000 rpm) were used to obtain deposits. Furthermore, composition of plating baths, such as the concentrations of samarium sulfamate as metal ions, glycine as complexer, and ammonium sulfamate, ammonium chloride and potassium chloride as conducting salts will be varied for different concentrations to examine their effects on electrodeposited cobalt-samarium alloys. The concentration of cobalt sulfate was constant at 0.05M. Plating baths used to obtain deposits are showed in Table 20:

TABLE 20
Plating baths of cobalt-samarium alloys
ItemBathSm sulfamateCo sulfateGlycineNH4 SulfamateNH4ClKClpH
111 M0.05 M0.15 M5.7
261 M0.05 M5.9
391 M0.05 M0.15 M5.0
4101 M0.05 M0.15 M4.0
5111 M0.05 M0.15 M3.0
6121 M0.05 M0.15 M2.0
7130.75 M  0.05 M0.15 M5.8
8140.5 M  0.05 M0.15 M5.9
9150.25 M  0.05 M0.15 M5.9
10161 M0.05 M0.05 M5.9
11171 M0.05 M0.10 M5.8
12181 M0.05 M0.20 M5.7
13191 M0.05 M0.50 M5.5
1481 M0.05 M0.15 M  1 M5.2
15201 M0.05 M0.15 M0.75 M5.3
16211 M0.05 M0.15 M 0.5 M5.4
17221 M0.05 M0.15 M0.25 M5.4
18231 M0.05 M0.15 M1 M5.5
19241 M0.05 M0.15 M1 M5.6
(The pH values of the plating baths were measured at 25° C.; the pH values of bath 9, 10, 11 and 12 were adjusted by sulfamic acid.)

Bath 12, 11, 10, 9 and 1 were used to study the effects of pH from 2 to 6; bath 15, 14, 13 and 1 were used to study the effects of the concentration of samarium sulfamate from 0.25 to 1M; bath 18, 19. 1, 16 and 17 were used to study the effects of the concentration of glycine from 0.05 to 0.5 M; bath 1, 22, 21, 20 and 8 were used to study the effects of the concentration of ammonium sulfamate from 0 to 1 M; bath 8, 23 and 24 were used to study the effects of different types of conducting salts in the plating baths.

d. Pretreatment and Post-Treatment

Before electrodeposition, the brass panels were mechanically cleaned, soaked in 0.1M NaOH solution for 10 min., rinsed in deionized water, immersed in 10% HCl for 30 seconds and than rinsed with deionized water. Unless otherwise noted, the total charge passed was 50 coulombs; solutions were not agitated during electrodeposition (except in the rotating disk electrode, RDE).

After the deposition of Co—Sm alloys for 50 coulombs, the deposits were removed from plating solution, rinsed with deionized water, and dried with nitrogen gas. Disk-shaped specimens of diameter of 6.4 mm (specimen area=31.7 mm2) were die-punched out from deposits for analysis.

e. Characterization and Analysis

The samarium

SmSm+Co(at%)

and cobalt deposit content

CoSm+Co(at%)

were determined by an energy dispersive x-ray spectroscopy (EDS) with a Kevex detector in a Cambridge SEM (see characterization section in Hull Cell studies); the mass of deposited cobalt was measured by a PerkinElmer flame atomic absorption spectrometer (AA, mode 631); the crystal structure, orientation, phase identification and grain size were determined by a PANalytical x-ray diffraction system (XRD, model X'Pert Pro) (see characterization section in Hull Cell studies); the surface morphology, microstructure and grain size were observed by a JEOL scanning electron microscopy (SEM, model JSM-6700F); magnetic properties were determined by a ADE Tech. vibrating sample magnetometer (VSM, model 1660). Unless otherwise noted, the experimental data presented in discussion sections are restricted to deposits with a metallic appearance.

f. Flame Atomic Absorption Spectroscopy (AA)

The specimen was dissolved by nitric acid than diluted with deionize water to 50 mL as the analytical solution for FAAS. By FAAS, the lowest cobalt concentration can be detected (limit of quantitation, LOQ) is 1 ppm; the concentration depart from linearity (limit of linearity, LOL) is 5 ppm. If the cobalt concentration of solution was out of applicable range (1-5 ppm), solution will be properly diluted or concentrated and the measurement was repeated. The Co concentration of the solution can be calculated from its absorption referring to the calibration curve. The mass of cobalt WCo in the specimen can be obtained by solution concentration and volume.

Calculation of the mass of Sm (W51,) and current efficiency (CE): The mass of Sm in the specimen Wsm can be calculated by:

WSm=WCoSm(wt%)Co(wt%)(Equation4)

where Wco was obtained by AA and the Sm and Co content were from EDS.

Calculation of the current efficiency (CE): The current efficiency (CE) can be calculated by the charge used to obtain metal (Sm and Co) divided by the charge (50 coulombs) passes during the electrodeposition. The charge used to obtain Sm and Co can be calculated by:

C=CCo+CSm=F(WCoZCo2+MCo+WSmZSm3+MSm) (unit:coulomb)(Equation5)

where C is the charge used to obtain metal (Co and Sm); Z is the valance of the metal ion (ZCO2+=2 and Zsm=3); F is Faraday's constant, the charge carried by a mole of electrons=96,500 C/mol; M is the atomic mass (MCo=58.93 g and MSm=150.36 g); W is the mass of elements in the specimen. The calculation of CE assumed the Sm and Co in deposit were obtained from charge transfer reactions (the reduction of Co2+ and Sm3+) rather than chemical reactions (the precipitation of Co(OH)2 and Sm(OH)2). The electrons passed during the electrodeposition were used only on the reduction of Sm, Co and H2. This experimental assumption provided a first approximation of CE in Co—Sm electrodeposition.

g. Scanning Electron Microscopy (SEM)

SEM provides specimen surface images by collecting the secondary or back scattering electrons emitted from specimen after the electron bombardment [P. J. Goodhew, J. Humphreys and R. Beanland, Electron microscopy and analysis (3rd edition), Taylor & Francis, (2001), pp. 196-205]. These images were used to study the topography, morphology and microstructure of electrodeposited Co—Sm alloys obtained from different solution and operating conditions.

h. Vibrating Sample Magnetometer (VSM)

VSM [D. Jiles, Introduction to magnetism and magnetic materials, Chapman & Hall, New York, (1991), pp. 47-53], a gradiometer measuring the magnetic induction difference with/without the specimen, gives a direct measurement of the magnetization (M) under applied magnetic field (H). A schematic of a typical VSM is shown in FIG. 62.

The disk shaped specimen was placed on a quartz sample holder and was vibrated with fixed frequency in C-C′ direction. The magnetic field was applied parallel (A-A′ direction) and perpendicular (B-B′ direction) to film plane to obtain in-plane and perpendicular magnetic properties, respectively. The magnetic field swept between ″10,000 and 10,000 Oe was used to obtain the hysteresis loop shown in FIG. 63.

Important magnetic properties of electrodeposited Co—Sm alloys can be obtained from its hysteresis loop as follows: He (Coercivity): the magnetic field needed to reduce the magnetic induction to zero after the material has been saturated (fully magnetized). Ms (Saturation Magnetization): maximum magnetization obtained in the hysteresis loop. Mr (Remanence): magnetization at applied magnetic filed equals to zero. Squareness (or the reduced remance): Mr/Ms. BHmax (Maximum Energy Product): the energy required to demagnetize a permanent magnet.

i. Effect of CD and Solution Temperature

Alloy Composition: Samarium deposit content (at %) increased with increasing current density (CD). CDmax (the highest CD to obtain metallic deposits) was extended by elevated solution temperatures (FIG. 64(a)). At 25° C., CDmax was limited to 50 mA/cm2 (Sm=14.5 at %), whereas for a solution temperature of 60° C., CDmax increased to 500 mA/cm2, resulting in deposit Sm content of 32.1 at %. Depending on CD and solution temperature, deposits of Sm content between 0 and 32 at % could be obtained from bath 1 (1 M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine) which satisfies the stoichiometric compositions of intermetallic Sm2Co17 (10.5 at %) and SmCo5 (16.7 at %) and, in addition, Sm2Co7 (22.2 at %) and SmCo3 (25 at %). Therefore, Co—Sm alloys of sufficient Sm content for Sm—Co magnets (i.e. Sm2Co17 and SmCo5) can be produced by electrodeposition.

Current efficiency (CE) calculated from the charge needed to obtain metals (i.e. Sm and Co) divided by the total charge passed in electrodeposition (50 coulombs). CE dropped sharply between 2 mA/cm2 (66%) and 100 mA/cm2 (20%) with little change between 100 mA/cm2 and 500 mA/cm2 in the 60° C. solution, whereas the CE decreased almost exponentially in the 25° C. electrolyte, as shown in FIG. 64(b). CEs were higher at elevated solution temperatures.

To show individual changes in reduced products in electrodepositions, a nomialized charge plot is provided in FIG. 65. Normalized charge or charge ratio indicates the electroreduction of a specified species. Calculation of normalized charges was based on two assumptions. First, all electrons supplied to the cathode contributed to the production of Sm, Co and H2 only; reduction of glycine or other species were ignored. Second, Sm and Co in deposits were obtained by electroreduction of reactants (ions) at the cathode.

Precipitates of Sm(OH)3 and Co(OH)2 were not considered. Normalized charges of Sm and Co were calculated from charges required to obtain the respective metals in the deposits divided by the total charge passed (50 coulombs). Normalized charge of H2 was obtained by difference (subtracting charges to produce Sm and Co from the total charge).

From the variation of normalized charges with CD, as shown in FIG. 7, at 60° C., increased CD (from 2 to 500 mA/cm2) resulted in a sharp increase in Sm from 0.01 to 0.09 and a decrease in Co from 0.65 to 0.13 leading to a substantial increase in Sm deposit content from 1.3 to 32.1 at % (FIG. 64(a)). At 25° C., increased Sm deposit content from 3 to 15 at % with increasing CD from 2 to 50 mA/cm2 was due to a sharp decrease in Co (normalized charge from 0.5 to 0.07) and a slight change in Sm from 0.02 to 0.03.

The cathode potential was measured at various CDs and solution temperatures with a saturated calomel electrode (SCE) serving as the reference electrode. In alloy electrodeposition, the cathode potential can affect the composition of deposits [N. Ibl, Surf Tech., 10, 81, (1980)]. In addition, the cathode potential may change the nucleation rate (grain size) [M. Paunovic and M. Schlesinger, Fundamentals of electrochemical deposition, John Wiley & Son, Inc., New York, (1998), pp. 107-121], microstructures [H. Seiter, H. Fischer, and L. Albert. Electrochim. Acta, 2, 97 (1960)], and phases and orientation [N. A. Pangarov, J. Electroanal. Chem., 9, 70, (1965); N. A. Pangarov and S. D. Vitkova, Electrochim. Acta, 11, 1733, (1966)] of deposits. These characteristics govern magnetic properties of deposits. FIG. 66(a) shows the polarization curves of Co—Sm electrodeposition from bath 1 at various solution temperatures.

Lower solution temperatures resulted in more negative cathode potentials in the electrodeposition of Co—Sm alloys. The dependence of Sm deposit content on cathode potentials is shown in FIG. 66(b). A higher Sm content was obtained at more negative cathode potentials. A linear relationship was found between Sm deposit content and the cathode potential, combining the effects of CD and solution temperatures.

It should be noted that the co-deposition of Sm and Co was observed at cathode potentials much less negative than the equilibrium potential of Sm (E°Sm/Sm2+=−2.65 V vs SCE) [W. M. Latimer, The oxidation states of the elements and their potentials in aqueous solution, Prentice Hall, N.Y., (1953), p 289]. This result indicates the co-deposition mechanism is more complex than the direct electrodeposition of both Co and Sm from their respective aqueous ionic forms.

j. Crystal Structure

The dependence of crystal structures on CD and solution temperature was determined by XRD. FIG. 67 and FIG. 68 show XRD results of deposits obtained from bath 1 at various CDs and at 25 and 60° C., respectively. It is noted that increased CD resulted in deposits changing from crystalline (or semi-crystalline) to non-crystalline; no diffraction peaks of Co, Sm or Co—Sm intermetallics were found between 10 and 50 mA/cm2 at 25° C. and Sm or Co—Sm intermetallics between 2 and 500 mA/cm2 at 60° C.; Co peaks disappeared at 60° C. and 500 mA/cm2. Crystal structures of electrodeposited Co—Sm crystallites were dominated by α-Co (hcp) phases; neither β-Co (fcc) nor Sm (rhombohedral) phases were found in deposits.

At 25° C. (FIG. 67), (10.0), (00.2), (10.1) and (11.0) peaks of α-Co and very weak (20.1) peak of SmCo5 (hexagonal) and (20.2) peak of Sm2Co17 (hexagonal) were observed at 2 mA/cm2 (Sm=3 at %). In addition, weak (11.0) peaks of Sm(OH)3 were found in the deposits at 25° C., but not at 60° C. (see FIG. 68).

At 60° C. (FIG. 68), a strong (00.2) peak of a-Co was observed at 2 mA/cm2 decreasing with further increased CD at 10 and 25 mA/cm2. Mixed orientations of (10.1), (11.0) and (10.0) peaks appeared from 10 (Sm 2.3 at %) to 100 mA/cm2 (Sm=9.2 at %). The (00.2) peak disappeared at 50 mA/cm2. Crystalline peaks were not seen at 500 mA/cm2 (Sm=32.1 at %).

Effect of solution temperature on deposit crystal structures at various CDs (2, 25 and 50 mA/cm2) are compared in FIG. 69. At higher CDs (25 and 50 mA/cm2), decreased solution temperatures changed deposits from crystalline to non-crystalline structures, similar to the effect of the increase of CD on deposit crystal structures (see FIG. 68 & FIG. 69). At low CD (2 mA/cm2), the decrease of solution temperature from 60 (Sm=1.3 at %) to 25° C. (Sm=3 at %) resulted in the decrease of α-Co (00.2) peak intensity, and other a-Co orientations (i.e. (11.0), (10.0) and (10.1)) were observed at 25 and 40° C. These orientations were also seen at 60° C., and both 10 and 25 mA/cm2 (FIG. 68).

According to these XRD results, the change in deposit orientation follow the same trend as varying CD or solution temperature and are related to Sm deposit content (or cathode potential). Thus, XRD results are organized by dependence of cc-Co orientation on Sm deposit content at various CDs and solution temperatures in Table 21. Higher solution temperatures required higher Sm content to produce non-crystalline deposits.

TABLE 21
Dependence of XRD patterns on Sm deposit content at various CDs and
solution
DepositCurrent Density (mA/cm2)
T (° C.)Properties2102550100200300500550
60hcp peaks(00.2) s(00.2) m(00.2) w(10.1) w(10.1) w(11.0) w(10.0) wnon-non-
(10.1) w(10.1) m(11.0) w(11.0) w(10.0) mcrystallinemetallic
(11.0) m(11.0) w(10.0) m(10.0) m
(10.0) m(10.0) m
Sm (at %)1.32.35.47.49.211.216.732.1
40hcp peaks(00.2) m(00.2) w(11.0) w(10.0) wnon-non-non-metallic
(10.1) w(10.1) w(10.0) wcrystallinecrystalline
(11.0) m(11.0) w
(10.0) m(10.0) w
Sm (at %)1.36.27.78.313.0 18.2
25hcp peaks(00.2) wnon-non-non-non-metallic
(10.1) wcrystallinecrystallinecrystalline
(11.0) s
(10.0) s
Sm (at %)3.06.59.514.5 
weak Sm(OH)3 (10-50 mA/cm2 at 25° C.), Sm2Co17 and SmCo3 (2 mA/cm2 at 60° C.) peaks were not included
s, m and w compared the intensity of peaks, “s” = strong, “m” = medium, “w” = weak
non-crystalline defined as no a-Co peaks was found in XRD

As shown in FIG. 66(b), Sm deposit content can be correlated to the cathode potential. Therefore, it is important to analyze the change in orientation with cathode potential. According to Pangarov's calculation for hcp lattice [N. A. Pangarov, J. Electroanal. Chem., 9, 70, (1965)] in electrodeposition, an [00.1] orientation should be expected at very low overpotentials, and with increase in overpotential, the following orientation should appear: [10.1], [11.0], [10.0] and [11.2]. Later on, the experimental results [N. A. Pangarov and S. D. Vitkova, Electrochim. Acta, 11, 1733, (1966)] of electrodeposited Co (sulfate bath, pH5, 1580° C., 10-200 mA/cm2) were provided. Electrodeposited Co obtained at low overpotential (high solution temperatures and low CDs) resulted in a pure [00.1] orientation. Medium overpotential (at low solution temperatures) leads to [11.0] and [10.0] as a mixed orientation. High overpotential, low solution temperatures and high CDs, had [10.0] orientation; [10.1] orientation was included with other orientation.

The dependence of orientation of electrodeposited Co—Sm alloys is similar to electrodeposited hcp Co as strongly related to CD, solution temperature, Sm content (or cathode potential) in agreement with Pangarov's calculation of the hcp lattice and his experimental results of hcp Co deposition.

k. Distortion of HCP Lattice

For more carefully studying the XRD patterns obtained at 60° C. from bath 1, Bragg angles (2ΘB of hcp-Co (00.2) and (10.0) peaks were found varying with Sm deposit content (FIG. 119). Increased Sm deposit constant resulted in the decrease of Bragg angle of (00.2) plane from 44.425° (point c) to 44.179° (point d) for Sm content increased from 1.3 to 2.3 at %. On the other hand, Bragg angle of (10.0) plane increased from 41.643° (point a) to 41.798° (point b) for Sm content increased from 2.3 to 16.7 at %.

Lattice constants of these deposits were calculated and compared to pure Co in Table 22. After comparing lattice constants of deposits of different Sm content, it was observed that lattice constant a (parallel to the basal plane of hcp structure) decreased with increased Sm deposit content, whereas lattice constant c (perpendicular to the basal plane of hcp structure) increased. Changed lattice constants implied the distortion of hcp Co lattice by adding Sm atoms into Co matrix. The radii of Co and Sm atoms are quite different; the atomic radius of Co is 1.25 A and Sm is 1.81 A [J. P. Schaffer, A. Saxena, S. D. Antolovich, T. H. Sanders, S. B. Warner, The science and design of engineering materials, McGraw-Hill Componies, Inc, New York, (1999), p 773 I° Centre d'information du cobalt, Cobalt monograph, 35, Rue Des Colonies, Brussels, Belgium, (1960), p75]. The size misfit

RSm-RCoRCo=45%

between Sm and Co could cause Co lattice distortion by adding Sm atoms. The distortion caused by addition of Sm tended to elongate Co lattice along the c-axis and compress along the basal plane. Furthermore, such a distortion of Co lattice may generate internal stress leading to microcracks in deposits. This will be discussed in the next section.

TABLE 22
Dependence of Bragg's angles (219) of (10.0) and (00.2) planes
and lattice constants of electrodeposited Co-Stn alloys
on Sm deposit content
Bragg's Angle (2⊖B)Lattice Constant
CDSm Content(deg.)(Å)
(mA/cm2)(at %)(10.0)(00.2)ac
Pure Co 41.595* 44.528* 2.50710 4.06910
21.341.64344.4252.5044.078
 102.344.1794.100
 507.441.6862.502
30016.741.7982.495
*Bragg angles of pure Co (hep) (10.0) and (00.2) peaks were calculated
from the lattice constants of pure Co10 by Bragg's law (1) and the
relationship between lattice constants and interplanar spacing of (hkl)
plane of hep structure (2)
2dhklsinθB=λ(1)1(dhkl)2=43(h2+hk+k2a2)+l2c2(2) where λ (CuKα) = 1.54184Å⊖e: Bragg's angle of (hkl) planedhkl: interplanar spacing of (hkl) plane

1. Morphologies and Microstructures

SEM pictures in FIG. 70 and FIG. 71 show the dependence of morphology and microstructure on CD at 25 and 60° C., respectively. FIG. 72 and FIG. 73 compare the effect of solution temperature at 25 and 50 mA/cm2, respectively. The increase in Sm content due to increased CD (FIG. 70 & FIG. 71) or decreased solution temperature (FIG. 72 & FIG. 73) resulted in more microcracks and smaller particle sizes.

Microstructures of crystalline deposits were fiber-shaped nano-rods. With increased Sm content, deposits changed to non-crystalline structures (XRD results) consisting of tiny roundish particles.

Microstructures of crystalline deposits of Co—Sm alloys are similar to electrodeposited Co. According to Cavallotti et al [P. L. Cavallotti, E. Galbiati and T. Chen, in Electroplating Engineering and waste Recycle, D. D. Snyder, U. Landau, R. Sard (Eds.), ECS Pub., Pennington, N.J., 130, (1983)], celluar electrodeposited Co was obtained from pH 6.5 sulfamate or sulfate solutions and changed to dendritic growth by increasing CD. Outgrowing basal planes with “needle-shaped crystalline particles” were found in deposits at 100 mA/cm2 and 50° C. In addition, they found that the crystallite size was mainly influenced by CD and solution temperature [P. L. Cavallotti, E. Galbiati and T. Chen, in Electroplating Engineering and waste Recycle, D. D. Snyder, U. Landau, R. Sard (Eds.), ECS Pub., Pennington, N.J., 130, (1983)]. Increasing solution temperature increased grain size from tens of nanometers to several hundred nanometers. Increasing CD, the grain size decreased. Similar results of the change in grain size with CD and solution temperature were observed in electrodeposited Co—Sm alloys.

Particle size changed with CD, solution temperature and Sm contents. Particle size was measured and presented in Table 23 and FIG. 74. The results provided average values of particle size to illustrate the tendency rather than exact values at operating conditions. Because most of the deposit microstructures were fiber-shaped nano-rods lying on the film plane, their lengths were much larger than widths and heights and were measured individually. Particle lengths (σL) and widths (σW) in the in-plane direction were determined by SEM; seven particles were sampled and measured from a SEM picture under an operating condition, and an average of particle size in the in-plane direction (σ) was the arithmetic average of the particle length and width. Particle thickness (σ) (perpendicular to film plane) was calculated by the Scherrer's equation according to FWHM of α-Co (10.0) peaks in XRD patterns and represent the particle size in the perpendicular direction.

TABLE 23
Dependence of particle site on Sm deposit content at various CDs and
solution temperatures
T Current Density (mA/cm2)
(° C.)Particle Size (nm)2550100300500
60Length σL57550045030063
Width σW2520181314
σ|| = (σL + σW)/230026023415838
Thickness, σ22181510
Sm (at %)5.47.49.216.732.1
Crystal Structurecrystallinenon-
crystalline
40Length σL44020595non-metallic
Width σW221816
σ|| = (σL + σW)/223111155
Thickness, σ2017
Sm (at %)7.78.313.0
Crystal Structurecrystallinenon-
crystalline
25Length σL8060non-metallic
Width σW1813
σ|| = (σL + σW)/24937
Thickness, σ
Sm (at %)9.514.5
Crystal Structurenon-crystalline
Length σL and width σW of particles (in-plane direction) are measured by SEM
Thickness, σof particles (perpendicular direction) are calculated by Scherrer's equation according to α-Co (10.0) peaks in XRD
— means no α-Co (10.0) peaks found in XRD pattern non-crystalline deposit)

Increased Sm deposit content leads to decrease in particle size, and the change was more significant in the in-plane direction mainly due to the reduction of the particle length. For deposits of similar Sm content, higher solution temperatures resulted in larger particle sizes in the in-plane direction (σ), whereas particle size in the perpendicular direction (σ) varied little with solution temperature. From the view point of nucleation and growth theory of electrocrystallisation, higher current density resulted in a higher nucleation rate [J. C. Puippe and F. Leaman, Theory and practice of pulse plating, American electroplaters and surface finishers society, Orlando, Fla., (1986), pp. 17-39] reducing the average distance between crystallites, therefore, a decrease in particle size can be expected.

m. Magnetic Properties

The most important characteristics governing the quality of electrodeposited hard magnetic films (i.e. coercivity Hc, saturation magnetization Ms and squareness Mr/Ms) are grain size, crystal structure and orientation and the presence of alloying elements [L. T. Romankiw and D. A. Thompson, in Magnetic properties of plated films in Properties of Electrodeposits: Their Measurements and Significance, Electrochemical Society, Princeton, N.J. (1975), pp 389-426]. Magnetic hysteresis loops of deposits obtained at various CDs and solution temperatures were measured by VSM for an applied magnetic field scanning between −10K and 10K Oe. In-plane (μ) and perpendicular (⊥) measurements represent the magnetic field applied parallel and perpendicular to the film plane, respectively. Magnetic properties of Hc, Ms and squareness were obtained from hysteresis loops.

FIG. 75 gives examples of hysteresis loops obtained at 25 and 60° C. and at various CDs. It was noted that magnetizations (Ms) were easier in the in-plane direction than the perpendicular direction indicating the easy-axis (EA) along the in-plane direction and the hard axis (HA) along the perpendicular direction. At 25 and 60° C., Ms were higher than Ms, and they approached each other sooner as magnetic field increased. Ms is used for the following discussion regarding approaching magnetization saturation. On the other hand, Hc⊥ were higher than Hc, and they got closer to each other as CD increased. At 25 and 60° C., Ms and Hc⊥ decreased as CD increased. At constant CD, Ms and He increased as solution temperature increased from 25 to 60° C. These results can be correlated to the alloy compositions and crystal structures of deposits. When deposits changed from crystalline to non-crystalline structures with increased Sm content by increased CD, magnetic properties of deposits appeared more isotropic where the Ms⊥ were closer to Ms and the Hc⊥ were closer to Hc Ms and Hc⊥ decreased with increased Sm deposit content. The dependence of Ms and He on alloy composition and deposit crystal structure were observed in these hysteresis loops.

To further quantify the dependence of magnetic properties on alloy composition and deposit characteristics, particle size, crystal structures, Hc, Ms, and squareness of deposits were correlated to Sm content in FIG. 76.

For the electrodeposited Co—Sm alloys before heat treatment, deposit characteristics (crystal structure and grain size) were strongly dependent on Sm deposit content. With increase in Sm deposit content, deposits changed from hcp Co crystallites to non-crystalline structure, and grain size decreased as shown in FIG. 76(a). Ms depended on alloy composition. Ms decreased linearly with increased Sm deposit content and was in agreement with sputtered films” (FIG. 76(b)). He sharply but changed little in the in-plane direction by increase in Sm content; Het approached Hc when deposits were of a non-crystalline structure (FIG. 76(c)). At 60° C., dependence of coercivities on Sm deposit content (FIG. 76(c)) can be correlated to crystal structure and particle size (FIG. 76(a)). Deposits obtained at 25 and 40° C. also followed the same trend, see FIG. 74.

Hoffman has shown that the coercivity of ferromagnetic thin films depends on crystallite size [H. Hoffman, IEEE Trans. Magnetics, 9, 17, (1973). Smaller crystallites result in decrease in coercivity. His prediction was later confirmed by the experimental results of electrodeposited Co films by Armyanov et al. [S. A. Armyanov and S. D. Vitkova, Phys. Status Solidi A, 26, 553, (1974)],[S. A. Armyanov and S. D. Vitkova, Surf Tech., 7, 319, (1978)] who found that coercivity increased with particle size between 20 and 400 nm. For electrodeposited Co—Sm alloys, the in-plane particle size ((σ=(σLw)/2) were larger than the perpendicular particle size (σ⊥) (FIG. 76(a)) because of fiber-shaped microstructures lying on the film plane (FIGS. 71(c), (f), (i) & (1)). For a fiber-shaped nano-rod lying on the film plane along the in-plane magnetic field (FIG. 77(b)), should consider the average particle size intersected by the in-plane magnetization, (σ⊥+σw)/2≈σ. Hc⊥ should consider the average particle size intersected by perpendicular magnetic field, (σLw)/2=σ. For those fiber-shaped nano-rods lying on the film plane but not along the in-plane magnetic field, the average particle sizes intersected by in-plane magnetization are between (σ⊥+σw)/2 and (σ⊥+σL)/2 depending on the angle between the fiber axis and in-plane magnetic field.

The sharp reduction in Hci with increased Sm content (FIG. 76(c)) can be correlated to the significant decrease in σ (FIG. 76(a)). On the other hand, σ⊥ decreased little by increased Sm content resulting in a small change in Hc. Larger σ than σ⊥ could explain higher Hc⊥ than Hc of these deposits. When deposits became non-crystalline consisting of tiny roundish particles (FIG. 70(i) and FIG. 71(o)), Hc⊥ was closer to Hc because of similar σ and σ⊥ values. The coercivity of electrodeposited Co—P alloys and Co metal also depend on crystal structure of deposits [K. Miller, M. Sydow and G. Dietz, Magn. Magn. Mater., 53, 269, (1985)]; increasing P content leads to a change from crystalline to non-crystalline deposits, and coercivity decreased significantly. For electrodeposited Co—Sm alloys, increased Sm content also resulted in deposits changing from crystalline to non-crystalline (FIG. 76(a)). This can cause the decreased coercivities.

The squareness ratio (Mr/Ms) of deposits provides the preference of magnetization direction. For example, magnetization direction is closer to the in-plane direction when in-plane squareness is higher. For ferromagnetic materials, the magnetization direction strongly depends on the minimization of the total magnetic energy. In the absence of an external magnetic field, magnetization direction is mainly controlled by the magnetocrystalline anisotropy energy [R. C. O'Handley, Modern magnetic materials, John Wiley & Son, Inc., New York, (2000), pp. 179-215] and demagnetization energy [R. L. Comstock, Introduction to magnetism and magnetic recording, John Wiley & Son, Inc., New York, (1999), pp. 24-28]. Minimization of magnetocrystalline anisotropy energy prefers to align magnetization along certain crystallographic directions. In hcp crystals, the magnetization direction prefers to align in the [00.1] direction [R. C. O'Handley, Modern magnetic materials, John Wiley & Son, Inc., New York, (2000), pp. 179-215]. To minimize demagnetization energy, magnetization prefers to lie along the long axis of a particle because demagnetization energy is proportional to the demagnetization factor which has the smallest value in the long axis direction of a particle [R. L. Comstock, Introduction to magnetism and magnetic recording, John Wiley & Son, Inc., New York, (1999), pp. 24-28]. For example, in a long cylinder particle, demagnetization factor along this axis is zero and perpendicular to the axis is 0.5.

As discussed in the previous section, fiber-shaped nano-rods were found to lie on deposit surfaces (long axis of particle aligned along the in-plane direction); particle size in the in-plane direction is larger than the perpendicular. Higher in-plane squareness than perpendicular can be explained by the alignment of magnetization direction (in absence of external field) along the long axis of particles to minimize demagnetization energy. On the other hand, the effect of magnetocrystalline anisotropy energy was not significant. (00.2) peaks were observed (see Table 21: 25° C.: 2 mA/cm2, 40° C.: 2 and 10 mA/cm2, 60° C.: 2, 10 and 25 mA/cm2) indicating the [00.1] orientation was along the perpendicular direction in these deposits. However, magnetization did not align along the perpendicular direction and resulted in higher perpendicular squareness.

Compared to the squareness of crystalline and non-crystalline sputtered SmCos deposits, crystalline deposits have higher squareness (0.3-0.8) than non-crystalline (−0.2) [C. Prados and G. C. I ladjipanayis. J. Appl. Phys., 83, 6253, (1998)]. Deposits changed from crystalline to non-crystalline structure (FIG. 76(a)) by increased Sm content indicating reduction in squareness. Magnetic properties of non-crystalline deposits exhibit isotropic over anisotropic (easy and hard axis caused by crystallographic structures no longer exist). Therefore, Hc and Hc⊥ (FIG. 76(c)) of a non-crystalline Co—Sm deposit were quite close. The in-plane and perpendicular Ms were also closer when deposits were non-crystalline (see FIG. 75(c), 25° C. or FIG. 75(i), 60° C.). On the other hand, the in-plane squareness was still higher than perpendicular (FIG. 76(d)) probably because the demagnetization direction is still aligned along the in-plane direction for the reduction of demagnetization energy. For non-crystalline deposits, particle size in the in-plane direction is still larger than in the perpendicular direction (see Table 21 and FIG. 74).

n. Effect of Solution pH

Alloy Composition: Solutions of various pHs were adjusted from bath 1 by sulfamic acid to study the effect of solution pH on Sm content and current efficiency as shown in FIG. 78.

For solution pH between 2 and 6, increased CD or decreased solution temperature resulted in higher Sm deposit content. At low CDs (10 & 50 mA/cm2) and 25° C., increasing solution pH increased Sm deposit content and then reached a constant. At 60° C., Sm content increased with increased solution pH and at 100 and 300 mA/cm2, Sm content reached maxima at pH 5 and 4. It is also important to point out that increasing CD results in increased rate of water reduction leading to increase in pH at cathode surface. The pH at the cathode surface is affected by both CD and solution pH. The results in FIG. 78(a), however, a plot of Sm content and solution pH. Estimation of the pH at the cathode surface may be made by calculation.

According to the mechanism proposed by Schwartz et al [M. Schwartz, N. V. Myung, and K. Nobe, J. Electrochem. Soc., 151, C468, (2004)], a heterodinuclear complex containing Sm and Co ions and glycine resulted in the co-deposition of Sm with Co. It is known that the glycine structure depends on solution pH and is considered as cationic (+H3N—CH2COOH ) at low pH, dipolar (+H3N —CH2COO) at medium pH (˜6) and anionic (H2N—CH2COO) at higher pH; glycine structures at different pH might alter its complexing ability with Sm and Co ions and change the co-deposition rate or Sm and Co.

Higher CEs were obtained at lower CDs and higher solution temperatures at various solution pH (FIG. 78(b)). CE varied little with solution pH compared to CD and solution temperature (FIG. D6(b)).

o. Crystal Structure

In the previous section, it was observed that crystalstructuress of electrodeposited CoSm alloys from bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine, pH 6) were dominated by a-Co (hcp) with some SmCos and Sm2Co17 crystallites. Weak Sm(OH)3 (11.0) peaks were found in deposits obtained at 25° C., but not in deposits at 60° C. Two interesting crystal structure problems were studied by varying the pH of bath 1. First, can low pH solutions eliminate Sm(OH)3 in deposits? Second, does the decreased pH result in the hcp→fcc transition which is often found in electrodeposited Co at low pH? The second problem is not only interesting in crystallography but also important for magnetic properties of deposits because magnetic properties of hcp and fcc Co are quite different. For example, hcp-Co has only one easy-axis [00.1] for magnetization, whereas fcc-Co has four easy axis of [111]. The energy of magnetocrystalline anisotropy of hcp-Co is higher than fcc-Co [Centre d'information du cobalt, Cobalt monograph, Brussels, Belgium, (1960), pp 95-100]. The Curie temperature of hcp-Co is 887° C. and fcc-Co is 1121° C.

XRD results of deposits obtained at various solution pHs and. CDs are shown in FIG. 79 and FIG. 80 (25° C.), and FIG. 81 to FIG. 84 (60° C.). These results are summarized in Table 24. Most deposits obtained at 25° C. (10 and 50 mA/cm2) were non-crystalline except for pH 2 at 10 mA/cm2 containing (10.0) and (00.2) peaks of a-Co. Decreased solution pH did not eliminate Sm(OH)3; (11.0) peaks of Sm(OH)3 were found in deposits at 25° C. between pH 2 and 6. On the other hand, the Sm(OH)3 peak was not observed at 60° C. At 60° C. for different pHs, the change in orientation with increasing CDs were similar.

TABLE 24
Dependence of XRD patterns on Sm deposit content at various
CDs and solution pH
Solu-
TtionDepositCurrent Density (mA/cm2)
(° C.)pHProperties1050100300
256hcp peaksnon-crystallinenon-metallic
Sm (at %)2.37.4
5hcp peaksnon-crystallinenon-metallic
Sm (at %)1.77.0
4hcp peaksnon-crystallinenon-metallic
Sm (at %)1.73.1
3hcp peaksnon-crystallinenon-metallic
Sm (at %)0.72.7
2hcp peaks(00.2) mnon-non-metallic
(10.0) mcrystalline
Sm (at %)0.51.5
606hcp peaks(00.2) m(10.1) w(10.1) w(10.0) w
(10.1) w(11.0) w(11.0) w
(11.0) m(10.0) m(10.0) m
(10.0) m
Sm (at %)2.37.49.216.7
5hcp peaks(00.2) m(11.0) w(10.1) wnon-
(10.1) m(10.0) m(11.0) wcrystalline
(11.0) m(10.0) w
(10.0) s
Sm (at %)1.77.011.8 22.6
4hcp peaks(00.2) w(00.2) w(10.0) wnon-
(10.1) w(10.1) wcrystalline
(11.0) w(11.0) w
(10.0) m(10.0) m
Sm (at %)1.73.110.8 28.1
3hcp peaks(00.2) w(00.2) m(10.0) w(10.0) w
(10.1) w(10.1) w
(11.0) w(11.0) m
(10.0) m(10.0) m
Sm (at %)0.72.78.524.2
2hcp peaks(00.2) m(00.2) w(00.2) wnon-
(10.1) m(10.1) w(10.1) wcrystalline
(11.0) s(11.0) m(11.0) w
(10.0) s(10.0) m(10.0) w
Sm (at %)0.51.50.918.9
weak Sm(OH)3 (10~50 mA/cm2 at 25° C.), Sm2Co17 and SmCo3 (2 mA/cm2 at 60° C.) peaks were not included
s, m and w compared the intensity of peaks, “s” = strong, “m” = medium, “w” = weak
non-crystalline defined as no α-Co peaks was found in XRD

Generally, the α-Co (hcp) phase is stable at temperature below 417° C., and the {tilde over (β)}Co(fcc) phase is thermodynamically stable only above 417° C. However, it was reported that both phases can be obtained in electrodeposited Co with α-Co obtained at pH higher than 2.9 and β-Co at pH less than 2.4. (bath: Co sulfate, NaCl, 18° C., 12 mA/cm2) [J. Goddard and J. G. Wright, Brit. J. Appl. Phys., 15, 807, (1964)]. Nakahara et al. [S. Nakahara and S. Mahajan, Electrochem. Soc., 127, 283, (1980)] (bath: Co sulfate, NaCl and boric acid, 25° C., 10 mA/cm2) further proposed that the formation of metastable Co hydride at low pH could be the reason for the formation of β-Co. High density inclusions of Co(OH)2 were observed in high pH (˜5.7) deposits [S. Nakahara and S. Mahajan, Electrochem. Soc., 127, 283, (1980)]. It was concluded that the solution pH is the most important parameter governing the crystal structure of electrodeposited Co where deposits at high pH results in α-Co and low pH generates β-Co [S. Nakahara and S. Mahajan, Electrochem. Soc., 127, 283, (1980)].

For Co—Sm deposits from bath 1, β-Co crystallites were not found from solution pH from 6 to 2 at 10 & 50 mA/cm2 and 25° C. (FIG. 79 and FIG. 80), and at 10-300 mA/cm2 and 60° C. (FIG. 81 to FIG. 84)). Furthermore, inclusions of Co(OH)2 was not observed even for deposits obtained at pH 6. In Hull Cell studies, it was observed that in the absence of glycine, Co(OH)2 and Sm2O3.CoO mixtures were found in deposits. It has been reported that the Co-glycine complex can inhibit the formation of Co(OH)2 [C. F. Diven, F. Wang, A. M. Abukhdeir, W. Salah, B. T. Layden, C. F. Geraldes, and D. M. Freitas, Inorg. Chem., 42, 2774, (2003)]. The appearance of glycine probably prevents inclusions of Co(OH)2 by complexing with Co ions.

p. Morphologies and Microstructures

SEM pictures of deposits obtained at various solution pHs and CDs at 25° C. and 60° C. are shown in FIG. 65 and FIG. 66 and FIG. 67 and FIG. 590, respectively. Dependence of particle size on Sm deposit content at various pHs, temperatures and CDs is shown in FIG. 59.

At 25° C. (FIG. 85), there were more microcracks in deposits obtained at higher CDs or solution pHs. These changes resulted in high Sm deposit contents. At 60° C. (FIG. 87), microcracks were less significant compared to 25° C. even for high Sm deposit contents. For electrodeposited Co-rich Co—Sm alloys (before the formation of intermetallic compounds) the addition of Sm into the Co matrix probably induced internal stress by the misfit of lattice constants of Co and Sm. Further addition of Sm increased internal stress increasing microcracks in the deposits. Stacking faults and defects formed during the electrodeposition also resulted in internal stress. Elevated solution temperatures resulted in adsorbed atoms of higher mobility on the cathode surface facilitating their reaching kink or terrace sites that reduce defects and internal stress in deposits.

High magnitude (50,000×) SEM results at 25° C. show ridge-shaped microstructures were observed at pH 2 below Sm=3.1 at % (FIGS. 86(g) & (h)). Ridge-shaped microstructures were also observed in Co-rich Co-Fe28 and Ni—Co29 alloys. At pH 4 and 10 mA/cm2, tiny roundish particles mixed with large plate-shaped particles resulted in a wider distribution of particle size (FIG. 86(d)).

At 60° C., microstructures at pH 2 were more compact (FIGS. 88(g) & (h)). At 50 mA/cm microstructures were large ridge-shaped (FIG. 88(h)). With decreased CD to 10 mA/cm2, tiny protrusions embedded in the ridge-shaped matrix had a two-phase structure (FIG. 88(g)). For deposits with similar Sm deposit contents obtained at different pH baths could have quite different microstructure shapes and particle sizes (for example, FIG. 88(d) vs. (h), FIG. 86(a) vs. (i)).

Vicenzo and Cavallotti studied the growth modes of electrodeposited Co from sulfamate baths. Different pH baths led to different morphologies and microstructures. Three basic modes were identified as: outgrowth, lateral growth and cluster growth which were strongly dependent on solution pH. Increased particle size was found as decreased [A. Vicenzo, P. L. Cavallotti. Electrochim. Acta, 49, 4079, (2004)]. In Co—Sm electrodeposition, deposit microctructure varied with solution pH. However, the dependence of particle size on solution pH was not significant.

At a fixed solution pH, increased CD or decreased solution temperature generally brought about an increase in Sm deposit content leading to the reduction of particle size (FIG. 89). However, at pH 2 and 60° C. particle size increased by increasing CD from 10 to 50 mA/cm2 (FIG. 88). The dependence of particle size at various pH on Sm deposit content was more scattered compared to a fixed pH (pH6, FIG. 74). Different shapes of microstructures obtained at different pH make this dependence more complex.

q. Magnetic Properties

FIG. 90 gives magnetic properties of deposits obtained at 25, 60° C. and at various solution pHs. Ms values were dependent on alloy compositions but not on solution pH. Ms decreased with increased Sm deposit content and were in good agreement with sputtered deposits.

Similar to the results on the effect of CD and solution temperatures in FIG. 18(c), Hc decreased with increased Sm deposit content (FIGS. 90(c) & (d)) due to decreased (FIG. 89). Hc varied little with Sm deposit content (FIGS. 90(c) & (d)) for small change in σ with increased Sm content (FIG. 89) (see FIG. 77).

In-plane squareness was higher than perpendicular due to the alignment of magnetization direction along the in-plane direction to reduce the demagnetization energy. Squareness decreased with increased Sm content by the change from crystalline to noncrystalline structure.

r. Effect of Fluid Dynamics

Alloy Composition: The rotating disk electrode (RDE), which changes the mass transfer rate of electrolyte from the bulk solution to the cathode surface by varying the rotation rate, was used to study the effect of fluid dynamics on Co—Sm alloy deposition. Because the brass substrate was placed facing down to the button of the cell, when the RDE was stationary, hydrogen bubbles generated from water reduction accumulated on the cathode surface resulting in burnt deposits with poor adhesion (films fell off the electrode after electrodeposition). For this reason, deposits obtained from parallel electrodes (without agitation) were used in place of 0 rpm (deposits) in the following discussions to assess the difference with/without agitation.

Deposit Sm content sharply decreased, then reached a constant by increase in rotating rate (FIG. 91(a)). On the other hand, CE increased with increasing rotating rate (FIG. 91(b)). Agitation by RDE resulted in metallic deposits at 100 mA/cm2 in contrast to deposits obtained without agitation (parallel electrodes), but higher Sm deposit content were not obtained. Even though the concentration of Co ions (0.05M) was much lower than Sm ions (1M) in bath 1, the deposition rate of Co was much greater than Sm (FIG. 92). A higher rotating rate did not significantly increase the deposition rate of Sm; but enhanced substantially the deposition of Co and suppressed H2 gas evolution. This indicates that the decrease in Sm deposit content by a greater agitation rate (FIG. 91(a)) was due to the increase in Co deposition rate (Sm deposition rate remained unchanged). These results indicate mass transfer effects in the co-deposition of Co.

As discussed in the previous section on parallel electrodes,, the deposits were not metallic at 100mA/cm2. White powder and burnt regions appeared on deposit surfaces. These non-metallic regions were confirmed as mixtures of oxides and hydroxides in the Hull cell study. SEM pictures of the deposit at 100 mA/cm2 (parallel electrodes, no agitation) are shown in FIGS. 93(a)-(c). High density microcracks were found in the deposit (FIG. 93(a)) and the white powder had a porous microstructure (FIG. 93(c)). For deposits obtained with agitation (1000 and 2000 rpm), metallic deposits were obtained and microcrack densities remained the same (comparing FIGS. 35(a), (d) & (e)) but the porous microstructure disappeared.

s. Magnetic Properties

The morphology and microstructure of the non-metallic deposit (confirmed as mixtures of hydroxides and oxides) obtained at 100 mA/m2 by parallel electrodes was compared to metallic deposits in the previous section. The appearance of oxide and hydroxide phases in this deposit also caused the degradation of magnetic properties, such as Ms. Compared to to sputtered Co—Sm alloys with similar Sm content, this deposit (parallel electrodes, 25° C., no agitation, 100 mA/cm, bath 1) had much lower Ms (about 30% f sputtered) (FIG. 94(a)).

Similar to the results of magnetic properties discussed in previous sections, Hc and squareness of deposits showed strong dependence on Sm deposit content. Perpendicular He decreased significantly but in-plane Hc decreased little by increasing Sm deposit content. In addition, Hc was higher in the perpendicular direction. In-plane squareness was higher than perpendicular squareness; squareness decreased more in in-plane direction than in the perpendicular direction with decreased Sm deposit content.

t. Effect of Sm Sulfamate

Alloy Composition: Baths containing 0.05M Co sulfate, 0.15M glycine and Sm sulfamate varying from 0.25 to 1M were used to study the effect of Sm sulfamate on deposit properties. The decrease of Sm sulfamate concentration resulted in non-metallic deposits (Table 25). White powders consisting of Sm(OH)3 and Co(OH)2 (XRD) appeared in deposits obtained for Sm sulfamate concentration <1M at 60° C. and low CDs (25 & 50 mA/cm2). CDmax (the highest CD with metallic deposits) decreased as Sm sulfamate concentration decreased (Table 25).

In general, decreased Sm sulfamate concentration increased Sm deposit content (FIG. 95(a)). CE increased with decreased Sm sulfamate concentration (FIG. 95(b)). The deposition rate of Sm and Co was enhanced by decrease in Sm sulfamate concentration, whereas Hi evolution was suppressed (FIG. 96). However, by decrease in Sm sulfamate concentrations, Sm(OH)3 or Sm oxide was found in metallic appearing deposits (Table 25) at 25° C. (FIG. 97 & FIG. 98) and 60° C. (FIG. 99 & FIG. 100). The precipitation of Sm(OH)3 or Sm oxide in deposits may contribute to the increase in Sm deposit content with decreased Sm sulfamate concentrations.

TABLE 25
Deposit appearance obtained at various Sm sulfamate
concentrations, CDs and temperatures
[Sm sulfamate]
1M0.75M0.5M0.25M
25° C.
CD (mA/cm2)
 25MMMM
 50MMMb
100bbbB
60° C.
CD (mA/cm2)
 25Mwww
 50Mmww
100MMMb
300MMbB
500MbBB

u. Crystal Structures

XRD results in the previous section show weak (11.0) peaks of Sm(OH)3 in deposits obtained at 25° C. (FIG. 67). At 25° C. and 25 mA/cm2, similar XRD patterns were found for deposits obtained from baths containing various Sm sulfamate concentrations (FIG. 97). On the other hand, at 50 mA/cm2 (25° C., FIG. 98), the intensity of Sm(OH)3 (11.0) peaks became stronger when Sm sulfamate concentrations were less than 0.75M. Meanwhile, Co(OH)2 (11.0) peaks appeared for Sm sulfamate concentrations less than 0.5M.

At 60° C., no hydroxide peaks were found in deposits from bath 1 (Sm sulfamate=1M) (FIG. 68). However, Sm(OH)3 (11.0) and Co(OH)2 (11.0) peaks appeared for Sm sulfamate concentrations of 0.75M (60° C. and 50 mA/cm2) or less with peak intensities increasing with decreasing Sm sulfamate concentrations (FIG. 99). A (111) peak of SmO was observed when Sm sulfamate concentration reached 0.25M. With further increase of CD to 100 mA/cm2 (FIG. 100), peaks of CoO and SmCoO3 were found.

Generally, decrease in Sm sulfamate concentrations resulted in more hydroxides (Sm(OH)3 and Co(OH)2), except for deposits obtained at 25° C. and 25 mA/cm2. SmO, CoO and SmCoO3 were observed at 60° C. and 100 mA/cm2 for Sm sulfamate concentrations of 0.5M and less. Co(OH)2 and CoO were not only found in non-metallic deposits but also in metallic deposits (FIG. 98(b), FIG. 99(c) and FIG. 100(b)) degrading their saturation magnitization.

v. Magnetic Properties

Magnetic properties of deposits obtained at various Sm sulfamate concentrations are shown in FIG. 101. Ms values of non-metallic deposits were ⅕ to ¼ that of sputtered deposits. Co(011)7 or CoO in these non-metallic deposits (FIG. 40(a) and FIG. 42(a)) degraded Ms by reducing the ferromagnetic phase (metallic Co) to non-ferromagnetic phases (Co(OH)2 or CoO). For metallic deposits containing Co(OH)2 or CoO obtained at 25° C., 50 mA/cm2, [Sm sulfamate]=0.5M (FIG. 40(b)) and at 60° C., 100 mA/cm2, [Sm sulfamate]=0.5M (FIG. 42(b)), Ms values were also much lower than sputtered deposits.

Similar to previous observations (FIG. 76(c)), He depended on Sm deposit content. Higher Sm deposit content resulted in lower Hc (compare FIG. 101(e) with (a), (f) with (b)); the change in perpendicular Hc was more significant than in-plane Hc.

Squareness followed similar trends as discussed (FIG. 76(d)): in-plane squareness was higher than perpendicular, and squareness decreased with increased Sm deposit content (compare FIG. 101(g) with (a), (h) with (b)).

w. Effect of Glycine

Alloy Composition: Baths consisting of 1M Sm sulfamate, 0.05M Co sulfate and glycine varied from 0.05 to 0,5M were used to study the effect of glycine on deposit properties. Dependence of Sm content and current efficiency on glycine is in FIG. 102.

At 25° C., low glycine concentrations resulted in non-metallic deposits. Metallic deposits were obtained when glycine concentration was higher than 0.1M at 25 mA/cm2 and 0.15 M at 50 mA/cm2. For metallic deposits obtained at 25° C., Sm deposit content decreased with increasing glycine concentration. At 60° C. and low CDs (25 and 50 mA/cm2), Sm deposit content increased, reached a maximum and then decreased with glycine increased from 0 to 0.5M; highest Sm contents were obtained at 0.15M glycine. With further increase in CD to 300 mA/cm2, the highest Sm content was obtained at 0.1M glycine, and metallic deposits were not observed at glycine concentration below 0.1M. Highest Sm contents were obtained at glycine concentration between 0.1M (glycine: Co2+=2:1) and 0.15M (glycine: Co2+=3:1). CE increased with increased glycine concentration at 25° C., whereas had no significant dependence on glycine concentration at 60° C.

According to the mechanism proposed by Schwartz et al [M. Schwartz, N. V. Myung, and K. Nobe, J. Electrochem. Soc., 151, C468, (2004)], a heterodinuclear complex containing Sm and Co ions and glycine resulted in the co-deposition of Sm with Co. Different glycine to metal ions ratio may change the complex composition (i.e. types and concentrations) in solutions affecting the reduction of Sm and Co and resulting in different Sm deposit contents.

x. Crystal Structure and Microstructure

Addition of glycine prevents the precipitation of Co(OH)2 and Sm(OH)3 as shown in the Hull cell results. Formation of Co-glycine complex has been reported to inhibit the formation of Co(OH)2 in aqueous solutions31 and could be the reason of preventing the precipitation of hydroxides in deposits. More studies by parallel electrode deposition will be discussed in this section. At 25° C. weak Sm(OH)3 (11.0) peaks were found in deposits were found in deposits obtained from 0.15M glycine solutions (FIG. 103(b) or FIG. 67). However, at 60° C. 0.15M glycine effectively prevented the precipitation of hydroxides in deposits (FIG. 104(b)).

On the other hand, deposits obtained at 0.05M and 0.5M glycine concentrations had stronger Sm(OH)3 (11.0) peaks compared to 0,15M glycine (25° C.: FIG. 103, 60° C.: FIG. 104). Co(OH)2 (11.0) peaks were also found in these deposits (0.05 and 0.5M glycine). Adding too little (glycine: Co2+=1:1) or too much glycine (glycine: Co2+=10:1) to the solution did not prevent the formation of hydroxides (Co(OH)2 and Sm(OH)3).

y. Magnetic Properties

FIG. 105 shows the magnetic properties of deposits obtained at various glycine concentrations. For metallic deposits obtained from the solutions of glycine concentrations ranging from 0.05 to 0.5M had Ms values comparable to sputtered deposits (FIGS. 105(c) & (d)). Even in the presence of some hydroxides in deposits obtained at 0.05M and 0.5M glycine (FIG. 103 and FIG. 104), significant degradation of Ms values were not observed compared to the deposits containing hydroxides and oxides obtained at low Sm sulfamate concentrations (see FIG. 101).

In-plane and perpendicular He for deposits obtained by varying glycine concentrations did not change significantly (FIGS. 105(e)&(0), but perpendicular was greater than in-plane Hc. At 25 and 60° C., in-plane squareness was greater than perpendicular squareness.

z. Effect of NH4 Sulfamate

Alloy Composition: Schwartz et al [M. Schwartz, N. V. Myung, and K. Nobe, J. Electrochem. Soc., 151, C468, (2004)] studied the electrodeposition of Co—Sm alloys under agitation of the plating solution containing 0.9M Sm sulfamate, 0.12M Co sulfamate, 0.36M glycine and 0.9M NH4 sulfamate resulting in maximum Sm deposits of 8 at %. However, Hull cell results indicated that the presence of NH4 sulfamate in solution reduced Sm deposit content. In the absence of NH4 sulfamate (bath 1), a Sm deposit content of 26 at % was obtained at 60° C. and 650 mA/cm2 from an unagitated solution. A Sm deposit content of 32 at % was obtained with parallel electrodes from bath 1 at 60° C. and 500 mA/cm2 (unstirred).

Although addition of NH4 sulfamate resulted in decreased Sm content, it was of interest to study how NH4 sulfamate reduced Sm deposit content and affected deposit properties. Baths consisting of 1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine and NH4 sulfamate varied from 0 to 1M were used to study the effect of NH4 sulfamate. Addition of NH4 sulfamate resulted in decreased Sm content (FIG. 48(a)) confirming the Hull cell results. The decrease in Sm content was more substantial at higher CDs. CE increased with increased NH4 sulfamate concentration at 25° C. but varied little at 60° C. (FIG. 106(b)). Increased NH4 sulfamate concentration suppressed Sm deposition (FIG. 107(a)) and enhanced Co deposition (FIG. 107(b)) leading to decreased Sm content (FIG. 106(a)).

The heterodinuclear complex containing Sm and Co ions and glycine was proposed resulting in the co-deposition of Sm with Co [M. Schwartz, N. V. Myung, and K. Nobe, J. Electrochem. Soc., 151, C468, (2004)]. However, NH3 from the deprotonated NH4+ could compete with glycine to form other complexes with Co and Sm ions and result in changes in compositions of complexes in solution. This can decrease Sm deposit content by suppressing co-deposition of Sm.

aa. Crystal Structure

FIG. 108 and FIG. 51 compared the crystal structures of deposits obtained from solution with (bath 1)/without 1M NH4 sulfamate (bath 8) at 25 and 60° C., respectively.

At 25° C., non-crystalline deposits were obtained at CD over 10 mA/cm2 for both baths with/without 1M NH4 sulfamate. At 2 mA/cm2, only the α-Co (00.2) peak was found in the presence of 1M NH4 sulfamate (bath 8). In its absence (bath 1), addition of α-Co (10.1), (11.0) and (10.0), SmCo5 (hexagonal) (20.1) and Sm2Co17 (hexagonal) (20.2) peaks were observed. It was interesting to note the absence of the Sm(OH)3 peak in deposits from bath 8 for CD≧25 mA/cm2. At 60° C. (FIG. 109), the dependence of crystal structure on CD for deposits from bath 8 was similar to bath 1, except no (10.1) peak was observed. Furthermore, unlike deposits obtained at 25° C., 60° C. deposits did not exhibit the Sm(OH)3 peak.

bb. Morphologies and Microstructures

For deposits obtained at the same temperature and current density the presence of NH4 sulfamate resulted in much smaller microcracks in deposits than in its absence (FIG. 110). The former produced lower Sm deposit content leading to lower internal stressed deposits as discussed in the previous section.

Microstructures obtained from baths with/without 1M NH4 sulfamate were quite similar, except at 60° C. and 50 mA/cm2 (FIG. 111).

At 60° C. and 50 mA/cm2, rigid-shaped rods and spherical particles were observed in deposits obtained from bath 8 (with 1M NH4 sulfamate) (FIG. 111(b)). These microstructures were more compact compared to deposits from bath 1 (without NH4 sulfamate).

cc. Magnetic Properties

FIG. 112 shows the magnetic properties of deposits obtained from baths containing various NH4 sulfamate concentrations. Deposits obtained from solutions with/without NH4 sulfamate had similar Ms and were in accord with the Ms of sputtered deposits (FIGS. 112(a) & (b)). For the deposits (with NH4 sulfamate), Hc⊥ was much greater than Hc and decreased significantly by increasing Sm deposit content, whereas Hc varied little (FIGS. 112(c) & (d)). Both in-plane and perpendicular squareness decreased with increase in Sm content and the former significantly larger. The dependence of He and squareness on Sm content was not affected by the presence of NH4 sulfamate.

dd. Supporting Electrolytes

Alloy Composition: Addition of NH4 sulfamate in solution was found to suppress the deposition of Sm and enhance the deposition of Co resulting in decreased Sm content. As a supporting electrolyte, NH4 sulfamate could change or modify the glycinato-complex structure in solution thereby affecting the deposition of Sm and Co.

Taube and Gould [H. Taube and E. S. Gould, Acc. Chem. Res., 2, 321, (1969)] indicated that NH3 is not a good bridging ligand for electron transfer in redox reactions, and a metal ion-NH3 complex could result in low reaction rates. Further, inclusion of Cl— ions could accelerate redox reactions and is referred as a good bridging ligand or mediating group. Therefore, it was of interest to include the effect of bridging ligands on electrodeposition of Co—Sm alloys. Supporting electrolytes of 1M NH4 sulfamate, NH4Cl or KCl were added to bath 1 to study the influence of NH3 and Cl— bridging ligands on deposit composition and properties.

The highest Sm content of metallic deposits obtained from these baths ranked as: 1M KCl (18 at % at 50 mA/cm2) >no supporting electrolyte (14.5 at % at 50 mA/cm2)>1M NH4Cl (9.7 at % at 100 mA/cm2)>1M NH4 sulfamate (8.1 at % at 300 mA/cm2) from 25° C. solutions (FIG. 113(a)). Sm deposit content increased with the addition of Cl— and decreased with the addition of NH4+ (or NH3). Addition of KCl did not increase CDmax. Compared to no supporting electrolyte (50 mA/cm2), however, the addition of NH4Cl and NH4 sulfamate resulted in increased CDmax to 100 and 300 mA/cm2, respectively.

With increased solution temperatures (60° C.), addition of NH4Cl or NH4 sulfamate both extended the CDmax to 900 mA/cm2, whereas addition of KCl limited the CDmax to 300 mA/cm2 (FIG. 113(b)). The highest Sm deposit contents were: no supporting electrolyte (32 at % at 500 mA/cm2)>NH4Cl (27 at % at 900 mA/cm2) NH4 sulfamate (27 at % at 900 mA/cm2)>KCl (21 at % at 300 mA/cm2). Because of the lower limiting CDmax (300 mA/cm2), the highest Sm deposit content from the KCl-containing bath was not higher than deposits from NH4Cl or NH4 sulfamate baths. However, with CDs<300 mA/cm2, the Sm deposit content ranked: KCl>no supporting electrolyte NH4Cl>NH4 sulfamate, similar to the deposits from 25° C. baths. The dependence of Sm content on NH4+ and Cl— ions (as bridging ligands) was observed; at a fixed CD, the Sm contents obtained from Cl— containing solutions were higher than NH4-containing solutions. CEs decreased with increasing CD at 25° C.; CDs at 60° C. decreased but reached a minimum at 300 mA/cm2 and then increased slightly (FIG. 114).

At 25° C., Sm deposit content increased with increasing CD (FIG. 113(a)) mainly due to the decreased. Co deposition (FIG. 115(a)); the decrease of Co deposition was sharper for solutions in the absence of supporting electrolyte. At 60° C., increasing CD increased Sm deposition and decreased Co deposition (FIG. 115(b)). At both 25 and 60° C., Sm deposition was enhanced by addition of Cl— ions and suppressed by addition of NH4+ ions (FIGS. 115(c) & (d)).

ee. Morphologies and Microstructures

FIG. 116 and FIG. 117 show the SEM of deposit surface obtained from baths at temperature of 25 (at 25 mA/cm2) and 60° C. (at 300 mA/cm2), respectively. The morphology of deposits shows little effect of the supporting electrolyte.

ff. Magnetic Properties

FIG. 118 shows the magnetic properties of deposits obtained from solutions with/without supporting electrolytes.

Addition of NH4 sulfamate, NH4Cl or KCl did not degrade deposit magnetic saturation (Ms). These Ms values were close to sputtered deposits (FIGS. 118(a) & (b)). While Hc⊥ decreased sharply and Hc varied little with increased Sm content, the addition of various supporting electrolytes had little effect on deposit coercivity (FIGS. 118(c) & (d)). Both in-plane and perpendicular squareness (Mr/Ms) decreased with the former significantly higher than the latter. Addition of supporting electrolyte containing Cl—, especially KCl, resulted in higher in-plane squareness compared to deposits from other solutions at D60° C. (FIG. 118(1)).

6. Aqueous Electrodeposition of Magentic Co—Sm Alloys—Pulse Current (PC) Electrodeposition Studies

In pulse current (PC) electrodeposition studies, an interrupted cathodic current with square waveform is applied for a specific time period (Ton) and then returned to ground zero for another specific time period (Toff); such a pulse period consisting of Ton and Toff repeats during the electrodeposition. Three important features in PC electrodeposition are: peak current density (PCD), concentration relaxation of reactants and kinetic selected deposition [Ibl, J. C. Puippe and H. Angerer, Surf Tech., 6, 287, (1978); N. Ibl, Surf Tech., 10, 81, (1980).]. These characteristics of PC electrodeposition affect alloy compositions and crystal properties of deposits.

Pulse current results in a very high instantaneous peak current density and hence a very negative cathodic potential. Higher CDs or more negative cathodic potentials have been shown to increase Sm deposit content in DC electrodeposition studies. Hull cell studies also show this trend in PC electrodeposition. Therefore, PC electrodeposition of high peak current densities can increase Sm deposit content. In addition, a very negative cathodic potential also increases the nucleation rate and can change the particle size and microstructures of deposits.

Short Ton and longer Toff provides more relaxation of the reactant (metal ions or complexes) concentrations at the cathode surface preventing the depletion of reactants and minimizing mass transfer effects. This changes the alloy composition. OH— ions were generated only during Ton in PC electrodeposition (DC generates OH— ions continuously during electrodeposition.). This can prevent unwanted Sm and Co hydroxides/oxides in deposits.

Frequency of pulse current also change alloy compositions. Unlike DC obtaining deposits steadily during electrodeposition, PC enables the kinetic selected deposition by increased frequency. In alloy electrodeposition, higher frequency increases deposit content of metal with higher reduction rates.

Disclosed herein is how PC electrodeposition parameters affected alloy properties (i.e., composition, crystal and magnetic properties), and correlate deposit magnetic behavior to other alloy properties. In addition, deposits obtained by DC and PC electrodeposition are compared.

The main goals of PC electrodeposition include: Obtaining high Sm deposit content of Co—Sm alloys; Determining the dependence of the deposit properties of Co—Sm alloys on PC electrodeposition parameters; Studying the relation between the deposit magnetic properties and PC electrodeposition parameters; Comparing deposit properties obtained by DC and PC electrodeposition.

FIG. 120 shows the experimental flowchart of a PC experiment which mainly includes four parts: pretreatment of cathode, PC electrodeposition, post-treatment of specimen and characterization. Definitions, setup and design of PC electrodeposition, pretreatment and post-treatment, and the characterization and analysis of the specimens will be detailed in following sections.

a. Definitions and Parameters of PC Electrodeposition

PCD (peak current density) is the maximum CD in one complete pulse cycle. Ton, is the time duration of the on-current in one complete pulse cycle. Toff is the time duration of the off-current in one complete pulse cycle. Period is the total time duration in one complete pulse cycle, period=Ttotal=Ton+Toff. Frequency is defined as the number of complete cycles per second,

f=1period=1Ton+Toff.

Duty cycle (y) is defined as the ratio of Ton to period,

γ=TonTon+Toff=Ton·f.

PCDmax is defined as the highest PCD to obtain metallic appearing deposits.

b. Setup of PC Electrodeposition

FIG. 122 shows the setup of the PC electrodeposition system. A Kraft Dynatronixpower generator (model DRP 20-5-10) served as the power source for PC electrodeposition. A coulometer was used to measure the charge passed during electrodeposition. The deposits were obtained in a 250 ml beaker filled with the plating bath of 240 ml. Brass panels (2×1.9 cm) served as cathodes and a platinum sheet (3×36 cm) was used as the anode; the distance between the cathode and the anode is 4 cm. A shielding panel with a 2×2 cm opening window, designed by the simulation result of ANSYS (a finite element analysis software), was placed between the cathode and anode to provide a uniform current density distribution on cathode to minimize the thickness variation of the deposit.

C. Design of Experiments

In this study, operating conditions for PC electrodeposition, such as peak current density, solution temperature, duty cycle, frequency and Ton will be varied. Various peak current densities (100-1200 mA/cm2), bath temperatures (25-60° C.), duty cycle (0.001-0.3), frequency (10-2 k Hz), and Ton (0.05-2 ms) were used to obtain deposits. Bath 1 (1M samarium sulfamate, 0.05M cobalt sulfate, 0.15M glycine) was used to determine the key variables in the PC co-electrodeposition of Co—Sm alloys.

d. Pretreatment and Post-Treatment

Before electrodeposition, the brass panels were mechanically cleaned, soaked in alkaline 0.1M NaOH solution for 10 min., rinsed in deionized water, immersed in 10% HCl for 30 seconds and than rinsed with deionized water. Unless otherwise noted, the total charge passed was 50 coulombs; solutions were not agitated during electrodeposition.

After the deposition of Co—Sm alloys for 50 coulombs, the deposits were removed from plating solution, rinsed with deionized water, and dried with nitrogen gas. Disk-shaped specimens of diameter of 6.4 mm (specimen area=31.7 mm2) were die-punched out from deposits for analysis.

e. Characterization and Analysis

The samarium

SmSm+Co(at%)

and cobalt deposit content

CoSm+Co(at%)

were determined by an energy dispersive x-ray spectroscopy (EDS) with a Kevex detector in a Cambridge SEM; the mass of deposited cobalt was measured by a Perkin Elmer flame atomic absorption spectroscopy (AA, mode 631); the crystal structure, orientation, phase identification and grain size were determined by a PANalytical x-ray diffraction system (XRD, model X'Pert Pro); the surface morphology, microstructure and grain size were observed by a JEOL scanning electron microscopy (SEM, model JSM-6700F); magnetic properties were determined by a ADE Tech. vibrating sample magnetometer (VSM, model 1660). Unless otherwise noted, the experimental data presented are restricted to deposits with a metallic appearance.

f. Effect of PCD and Solution Temperature

Alloy Composition: Bath 1 (1M Sm sulfamate, 0.05M Co sulfate, 0.15M glycine) was used to study the effects of peak current density (PCD) and solution temperature on alloy properties. Ton was maintained constant as 0.1 ms and duty cycle y of 0.1. FIG. 123 compares the effect of PCD (or CD) and solution temperature on Sm deposit content and current efficiency in PC and DC electrodeposition. Similar to DC electrodeposition, increased PCD resulted in increased Sm deposit content. At 25° C., the PCDmax of 1050 mA/cm2 was much higher than the CDmax of 50 mA/cm2 resulting in a higher maximum. Sm deposit content by PC (20.3 at %) than by DC (14.5 at %). On the other hand, at 60° C. although the PCDmax (2100 mA/cm2) was higher than CDmax (500 mA/cm2), maximum Sm content by PC (11.6 at %) was lower than by DC (32.1 at %) due to smaller

SmcontentPCD(orCD)

of PC electrodeposition. PC electrodeposition at 60° C. resulted in a lower maximum Sm content (11.6 at %) compared to 25° C. (20.3 at %). On the other hand, DC electrodeposition showed the opposite result of higher maximum Sm content at 60° C. (32.1 at %) than 25° C. (14.5 at %). This confirms the Hull cell studies. Increased PCD led to decreased current efficiency, and elevated solution temperatures resulted in higher current efficiencies in PC electrodeposition.

At 60° C., PC reduced Sm deposition and enhanced Co deposition compared to DC electrodeposition (FIG. 124) resulting in lower Sm contents in PC electrodeposition.

g. Crystal Structures

Compared to DC electrodeposition, Sm(OH)3 was not found in deposits by PC electrodeposition (FIG. 125). DC generates OH— ions continuously during electrodeposition. On the other hand, OH— ions were generated only during Ton in PC electrodeposition. Therefore, PC electrodeposition resulted in lower OH— ion concentration at the cathode surface and minimizing the folination of Sm(OH)3 at both 25 and 60° C.

Deposits obtained at 25° C. (FIG. 125, left) appeared non-crystalline for PCD higher than 200 mA/cm2. Unlike DC electrodeposition, non-crystalline deposits were not found in PC electrodeposition at 60° C. (FIG. 125, right). All deposits obtained at 60° C. were hcp crystallites, even up to 2100 mA/cm2. The changes in orientation with increased PCD (or Sm content) were different than DC and not in agreement with Pangarov's prediction [N. A. Pangarov, J. Electroanal. Chem., 9, 70, (1965).]. (00.2), (10.1), (11.0) and (10.0) peaks of hcp Co were observed for deposits without following any role.

h. Morphology and Microstructures

FIG. 126 and FIG. 127 shows the SEM of deposits obtained at 25 and 60° C., respectively. At 25° C., increased PCD resulted in increased Sm deposit content and decreased particle size, similar to DC. Also, microstructures changed from fiber-shaped nano-rods to roundish particles, and microcracks increased with increased PCD.

However, for deposits obtained at 60° C., particle size decreased significantly by increased PCD from 100 to 300 mA/cm2 (FIGS. 127(c) & (f)) but changed little from 300 to 2100 mA/cm2 (FIGS. 127(f), (i) & (1)). Ridge-shaped microstructures were observed in the deposits obtained at 100 mA/cm2 (60° C.).

i. Magnetic Properties

FIG. 128 shows the hysteresis loops of deposits obtained at various PCDs and solution temperatures. Similar to deposits by DC electrodeposition, magnetization was easier in the in-plane direction than the perpendicular direction indicating the easy-axis (EA) along the in-plane direction and the hard axis (HA) along the perpendicular direction. In-plane magnetization (M) was higher than perpendicular magnetization (M⊥). On the other hand, Hc⊥ was higher than Hc. With increased PCD deposits changed from anisotropic to isotropic magnetic behavior at both 25 and 60° C. Such a change was more significant for deposits obtained at 60° C. Ms and He increased as solution temperature increased from 25 to 60° C.

Magnetic properties of deposits by DC and PC electrodeposition are compared in FIG. 129. (Magnetization was more complete in the in-plane than the perpendicular direction so that Msii is used to represent Ms in the following discussion.) The strong dependence of magnetic properties of deposits on Sm content was similar to DC electrodeposition. PC Ms decreased linearly with increased Sm deposit content, in agreement with sputtered films [H. S. Cho, J. R. Salem, A. J. Kellock and R. B. Beyers, IEEE Trans. Magnetics, 33, 2890, (1997).] (FIGS. 129(a) & (d)). Increased Sm content caused decreased H⊥ but Hc varied little (FIGS. 129 (b) & (e)). Hc⊥ declined and approached Hc with increased Sm content. S was higher than S⊥ confirming the aligning of the easy axis of magnetization along the in-plane direction (FIGS. 129(c) & (f)). Without wishing to be bound by theory, it is beleieved that decreased S and S⊥ with increased Sm content was due to the increased non-crystallinity of deposits.

j. Effect of Duty Cycle

Alloy Composition: Bath 1 was used to study the effects of duty cycle on alloys properties with Ton=0.1 ms, PCD=200 or 500 mA/cm2. An increase in duty cycle (y) increased Sm content linearly at 25° C. and parabolically at 60° C. (FIG. 130(a)). Increased Sm content was more significant at the lower solution temperature (25° C.) and higher PCD (500 mA/cm2). On the other hand, increased y resulted in exponentially decreasing current efficiencies. The decrease was more significant for deposit obtained at 25° C. than at 60° C.

Increased y did not enhance the deposition of Sm but considerably suppressed the deposition of Co (FIG. 131) leading to increased Sm deposit content (FIG. 130(a)). It was concluded from the RDE experimental results that mass transfer effects were greater for Co than Sm co-deposition. Increased y (decreased Toff) caused a lower Co concentration at the cathode surface because less Co ions recovered from bulk solution for shorter Toff, resulting in the decrease of Co deposition.

k. Crystal Structures, Morphologies and Microstructures

Increased γ resulted in increased Sm content and changed deposits from crystalline to non-crystalline structures at both 25 and 60° C. (FIG. 132). Increased γ also induced more microcracks in deposits at both 25° C. (FIGS. 133(a) & (d)) and 60° C. (FIGS. 133(g), (j) & (m)) and led to decreased particle size at 25° C. (FIGS. 133(c) & (f)) and 60° C. (FIGS. 133(i), (1) & (o)).

These deposit characteristics caused by increased Sm content were also observed for DC electrodeposition. Increased Sm content in the Co—Sm alloys could distort Co lattices probably changing the deposits from crystalline to non-crystalline structures and leading to more microcracks.

1. Magnetic Properties

FIG. 134 shows the magnetic properties of deposits obtained at various γ. Similar to the previous observation (effect of PCD and solution temperature on magnetic properties), magnetic properties of deposits obtained at various γ can be correlated to their Sm content which controlled the crystal structure and particle size. Ms values decreased with increased Sm content. Ms obtained from various γ (0.025-0.3) at both 25 and 60° C. were in agreement with sputtered films [H. S. Cho, J. R. Salem, A. J. Kellock and R. B. Beyers, IEEE Trans. Magnetics, 33, 2890, (1997).]. Hc⊥ was higher than Hc. Increased Sm content decreased Hc⊥ but Hc varied little (FIGS. 134(b) & (e)). Hc⊥ obtained at 25° C. decreased linearly with increasing Sm content, but at 60° C. Hc⊥ decreased gradually for Sm content less than 10 at % then dropped sharply. S was higher than S Both S⊥ and S⊥ decreased with increased Sm content (FIGS. 134(c) & (f)).

m. Effect of Frequency

Alloy Composition: Bath 1 was used to study the effect of frequency. Duty cycle γ was kept constant at 0.1, solution temperature at 25° C. and PCD at 100, 250 and 500 mA/cm2. Increased frequency resulted in linear decrease in Sm deposit content (FIG. 135(a)).

Increased frequency enhanced Co deposition, especially at low PCD, but Sm deposition varied little (FIG. 136) resulting in decreased Sm deposit content. These results indicate that Co deposition rate was greater than Sm deposition rate during electrodeposition of Co—Sm alloys leading to decreased Sm content at higher frequencies. Generally, increased frequency resulted in increased current efficiency, except for the deposits obtained at 100 mA/cm2 and 2000 Hz.

n. Crystal Structures and Morphologies

With increased frequency, deposits obtained at 100 Am/cm2 and 25° C. changed from non-crystalline to crystalline (FIG. 137 left) probably due to decreased Sm content. At low frequencies (100 Hz, 7.1 at % Sm), characteristic peaks for crystallites were not found indicating non-crystalline deposits. At medium frequencies (200-1 k Hz, 6.1-4.7 at % Sm), crystallites of mixed (10.0) and (11.0) peaks were observed. At high frequencies (2 kHz, 1.7 at % Sm), there was a strong (00.2) peak. Deposits obtained at 25° C. and frequencies between 100 and 1 kHz had similar morphologies (FIG. 137 right). Microcracks were present in these deposits. However, for the deposits obtained at 100 mA/cm2 and high frequency of 2 kHz, about 30% of the surface of the brass substrate was not covered by CoSm deposits (FIG. 137(d) right, FIGS. 138(a) & (b)). Increased PCD to 500 mA/cm2, the brass substrate was fully covered by deposits (FIGS. 138(c) & (d)). In other words, low coverage of deposits occurred only at high frequencies (2 kHz) and low PCD (100 mA/cm2).

o. Magnetic Properties

FIG. 139 shows the effect of frequency on magnetic properties in PC electrodeposition (25° C., γ=0.1). Similar to previous observations, magnetic properties of deposits were dependent on the SM content.

Ms values decreased linearly with increased Sm content and are in agreement with sputtered films [H. S. Cho, J.R. Salem, A. J. Kellock and R. B. Beyers, IEEE Trans. Magnetics, 33, 2890, (1997).] (FIG. 139(b)) in the frequency range between 100 and 2,000 Hz. Increased Sm content decreased Hc⊥ but Hc varied little (FIG. 139(c)). Although both S and S⊥ decreased with increased Sm content (FIG. 134 (d)), the decrease was less significant compared to the effect of PCD and solution temperature (FIG. 129) and duty cycle (FIG. 134).

p. Effect of Ton

Alloy Composition: Bath 1 was used to study the effects of Ton. Compared to studies discussed earlier, short Ton (0.1-2 ms) and long Toff (98-99.9 ms) were investigated to maintain the solution composition at cathode surface close to the bulk composition upon initiation of each pulse. By doing this, not only the effect of Ton but also the deposition rates of Sm and Co can be investigated. Pulse period (Ton+Toff) was fixed at 100 ms, solution temperature at 25° C. and PCD between 100 and 500 mA/cm2.

Increased Ton resulted in a parabolic increase in Sm content (FIG. 140(a)) at both 500 and 1000 mA/cm2. Metallic deposits were obtained at Ton below 2 and 1 ms for 500 and 1000 mA/cm2, respectively. Higher PCD reduced the maximum Ton for metallic deposits. The effects of Ton on individual Sm and Co deposition were quite different. Increased Ton led to a slightly higher Sm deposition rate (FIG. 141). On the other hand, Co deposition increased, reached a maximum and decreased with increased Ton. After reaching a maximum, the decrease in Co deposition with increased Ton confirmed the mass transfer effects observed in the rotating disk electrode results. CE increased, reached a maximum and decreased with increased Ton (FIG. 140(b)).

q. Crystal Structure and Microstructures

FIG. 142 show the XRD (left) and the SEM (right) results of deposits obtained at 25° C. and 1000 mA/cm2 with various Ton. The (11.0) peak of Sm(OH)3 was found in the deposit obtained at Ton, of 0.1 ms (FIG. 142(d)). With further increased Ton, Sm(OH)3 peaks disappeared. At short Ton of 0.1 ms, (00.2), (10.0) and (11.0) peaks of hcp-Co appeared in the deposit. When Ton increased to 1 ms, only the (10.0) peak was found. The change of deposit orientations with increased Ton is similar to DC electrodeposition at 60° C. Both of these changes were due to increased Sm deposit content. Increased Ton resulted in increased Sm content also leading to significant decrease in microstructure size (FIG. 142, right). Again, the dependence of crystal orientation and particle size on Sm content was observed.

r. Magnetic Properties

The effects of Ton on magnetic properties are shown in FIG. 143. As discussed in previous sections: Ms values depended on alloy composition and decreased linearly with increased Sm content in agreement with sputtered films. Hc⊥ was higher than Hc∥ and decreased with increasing Sm content (FIG. 143(c)). S were larger than S⊥ and both decreased with increased Sm content (FIG. 143(d)).

s. Deposition Rates of Sm and Co

In the previous section, long periods (100 ms), short Ton (0.1-2 ms) and low duty cycles (0.001-0.02) examined the effect of Ton on Sm deposit content. Because short Ton consumed less metal ions in a single pulse and low duty cycle provided greater relaxation time for the recovery of metal ion concentrations, the metal ion concentration at the cathode surface probably remain close to the bulk solution concentration just before the beginning of each pulse. Therefore, we assumed that at the initiation of each pulse the solution composition at the cathode equaled the bulk solution. Thus, the deposit contents at each pulse were identical, and Sm and Co in deposits were assumed the result of the reduction of Sm and Co ions to metals rather than to precipitation of Sm(OH)3 and Co(OH)2. This provided a first approximation in the calculation of the electrodeposition rates of Sm and Co.

The deposit content for each pulse was assumed to be identical. Therefore, the amount of electrodeposited Sm and Co per pulse at different Ton can be calculated, as shown in Table 26. Electrodeposition rates at different time can be obtained by taking the difference in deposit content (amount of Co and Sm) for pulses and divided by the difference in time duration.

TABLE 26
Calculation of reaction rates of Sm and Co
Pulse InformationDeposits for 50 CDeposits per PulseReaction Rate
Toncharge perpulseSmCoSmCoTimeSmCo
(ms)pulse (C)number(mole)(mole)(mole)(mole)(ms)(mole)/s(mole)/s
0.11.90E−042631582.3E−082.8E−068.7E−141.1E−110.058.7E−101.1E−07
0.35.70E−04877191.4E−075.0E−061.6E−125.7E−110.207.5E−092.3E−07
0.59.50E−04526322.9E−077.6E−065.6E−121.4E−100.402.0E−084.4E−07
0.71.33E−03375943.2E−077.1E−068.4E−121.9E−100.601.4E−082.2E−07
1.01.90E−03263163.4E−076.7E−061.3E−112.5E−100.851.5E−082.2E−07
1.32.47E−03202433.9E−076.6E−061.9E−113.3E−101.152.1E−082.4E−07
1.63.04E−03164474.2E−076.3E−062.5E−113.8E−101.452.1E−081.9E−07
23.80E−03131584.5E−076.0E−063.4E−114.5E−101.802.1E−081.8E−07
charge per pulse (C) = PCD × area × Ton = 0.5(A/cm2) × 3.8 (cm2) × 0.0001 (sec) = 1.90E−04
pulse number = total applied charge/charge per pulse = 50/1.90E−04 = 263158
Deposits for 50 C (total applied charge): Sm: 2.3E−08 (mole) and Co: 2.8E−06 (mole) from AA and EDS
Deposits per Pulse: Deposits for 50 C/pulse number, Sm = 2.3E−08/263158 = 8.7E−14 (mole) Co = 2.8E−06/263158 = 1.1E−11 (mole)
Reaction Rate (at 0.05 ms): Deposits per Pulse/Deposit duration, Sm = 8.7E−14/0.0001 = 8.7E−10 (mole/s) Co = 1.1E−11/0.0001 = 1.1E−07 (mole/s)
Reaction Rate (at 0.2 ms = (0.1 ms + 0.3 ms)/2); Sm = (1.6E−12-8.7E−14)/(0.0003-0.0001) = 7.5E−09 (mole/s) Co = (5.7E−11-1.1E−11)/(0.0003-0.0001) = 2.3E−07 (mole/s)

The deposition rates of Sm and Co at various deposition times are plotted in FIG. 144. Deposition rates of Sm increased linearly with deposition time. On the other hand, the deposition rate of Co increased, reached a maximum, and then decreased with increased deposition time. Higher PCD of 1000 mA/cm2 resulted in higher deposition rates of both Sm and Co compared to the lower PCD of 500 mA/cm2. Higher PCD caused Co deposition rates to reach a maximum in a shorter deposition time (0.3 ms for 1000 mA/cm2 and 0.4 ms for 500 mA/cm2). Co deposition rates were much higher than Sm deposition rates (about 10-120 times higher depending on PCD and deposition time) indicating that Co deposition is faster than Sm in agreement with results on frequency effects. The decrease in Co deposition rates after the maximum indicates mass transfer effects in the electrodeposition of Co—Sm alloys consistent with the results of rotating disk electrode studies and the effect of Ton.

It will be apparent to those skilled in the art that various modifications and variations can be made in the present invention without departing from the scope or spirit of the invention. Other aspects of the invention will be apparent to those skilled in the art from consideration of the specification and practice of the invention disclosed herein. It is intended that the specification and examples be considered as exemplary only, with a true scope and spirit of the invention being indicated by the following claims.